Structure/Property Relationships in Irons and Steels

Structure/Property Relationships in Irons and Steels

Bruce L. Bramfitt, Homer Research Laboratories, Bethlehem Steel Corporation

Basis o f Mater ia l Selection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Role o f Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155

Ferrite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 156 Pearlite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 158 Ferrite-Pearl ite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 160 Bainite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 162 Martensite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 164 Austenite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 169 Ferrite-Cementite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 170 Ferrite-Martensite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171 Ferrite-Austenite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171 Graphi te . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 172 Cementi te . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 172

This Section was adapted from Materials 5election and Design, Volume 20, ASM Handbook, 1997, pages 357-382. Addi t ional information can also be found in the Sections on cast irons and steels wh ich immediately fo l l ow in this Handbook and by consult ing the index.

THE PROPERTIES of irons and steels are linked to the chemical composition, processing path, and resulting microstructure of the material; this correspondence has been known since the early part of the twentieth century. For a particular iron and steel composition, most properties depend on microstructure. These properties are called

structure-sensitive properties, for example, yield strength and hardness. The structure-insensitive properties, for example, electrical conductivity, are not discussed in this Section. Processing is a means to develop and control microstructure, for example, hot rolling, quenching, and so forth. In this Section, the role of these factors is described

in both theoretical and practical terms, with par- t icular focus on the role of microstructure.

Basis of Material Selection In order to select a material for a part icular

component , the designer mus t have an int imate

" "o" - grade 50). 2% nital + 4% picral etch. 200x Fig. :2 Microstructu r e p e a r l i t e interlamellar°f a typicalspacing.fUllY2%pearlitiCnital + 4%rail steelpicralShowingetch. 500xthe characteristic fine

Metals Handbook Desk Edition, Second EditionJ.R. Davis, Editor, p 153-173

Copyright © 1998 ASM International® All rights reserved.

www.asminternational.org

154/Structure/Property Relationships in Irons and Steels

k n o w l e d g e of w h a t p r o p e r t i e s are r equ i red . Con- s i d e r a t i o n m u s t be g i v e n to the e n v i r o n m e n t ( c o r r o s i v e , h i g h t e m p e r a t u r e , etc.) and how the c o m p o n e n t w i l l be f ab r i ca t ed (we lded , bo l ted , etc.) . O n c e these p rope r ty r e q u i r e m e n t s are es- t a b l i s h e d the m a t e r i a l s e l e c t i o n p r o c e s s can be- g in . S o m e o f the p r o p e r t i e s to be c o n s i d e r e d are :

Mechanical properties Other properties/ Strength characteristics

Tensile strength (ultimate Formability strength) Dmwability

Yield strength Stretchability Compressive strength Bendability

Hardness Wear resistance Toughness Abrasion resistance

Notch toughness Galling resistance Fracture toughness Sliding wear resistance

Ductility Adhesive wear resistance Total elongation Machinability Reduction in area Weldability

Fatigue resistance

Table 1 l is ts mechanica l proper t ies of se lec ted s t ee l s in va r ious hea t - t r e a t ed or c o l d - w o r k e d cond i t i ons .

In the s e l e c t i o n p roces s , wha t is r e q u i r e d for one a p p l i c a t i o n may be to t a l ly i n a p p r o p r i a t e for a n o t h e r app l i ca t i on . For e x a m p l e , s tee l b e a m s for a r a i l w a y b r idge r equ i r e a t o t a l l y d i f f e r en t se t o f p rope r t i e s than the s tee l r a i l s that are a t t ached to the wooden t ies on the b r i d g e deck. In d e s i g n i n g the b r idge , the s tee l mus t h a v e su f f i c i en t s t r eng th to w i t h s t a n d s u b s t a n t i a l app l i ed loads . In fact , the d e s i g n e r wi l l g e n e r a l l y s e l ec t a s t ee l w i t h h ighe r s t r eng th than a c t u a l l y r equ i red . A l so , the d e s i g n e r k n o w s that the s t ee l m u s t have f r ac tu re t o u g h n e s s to res i s t the g r o w t h and p r o p a g a t i o n of c r acks and m u s t be c a p a b l e of b e i n g w e l d e d so tha t s t ruc tu ra l m e m b e r s can be j o i n e d w i t h o u t s a c r i f i c i n g s t r e n g t h and toughness . The s t ee l b r i d g e m u s t a l so be co r ro s ion res i s tan t . Th i s can be p r o v i d e d by a p r o t e c t i v e l aye r of pa in t . I f p a i n t i n g is not a l l owed , s m a l l amoun t s o f ce r t a in a l l o y i n g e l e m e n t s such as c o p p e r and c h r o m i u m can be added to the s tee l to i nh ib i t or r e d u c e c o r r o s i o n ra tes . Thus , the s t ee l s e l ec t ed for the b r idge wou ld be a h i g h - s t r e n g t h l o w - a l l o y ( H S L A ) s t ruc tu ra l s tee l such as A S T M A572 , g r a d e 50 or p o s s i b l y a w e a t h e r i n g s t ee l such as A S T M A588. A t) ;pical H S L A s tee l has a fe r r i t e - p e a r l i t e m i c r o s t r u c t u r e as seen in Fig. 1 and is m i c r o a l l o y e d wi th v a n a d i u m and /o r n i o b i u m for s t r e n g t h e n i n g . (Microalloying is a t e r m used to d e s c r i b e the p roces s of u s i n g s m a l l a d d i t i o n s of c a r b o n i t r i d e f o r m i n g e l e m e n t s - - t i t a n i u m , vana- d ium, and n i o b i u m - - t o s t r e n g t h e n s t ee l s by g ra in r e f i n e m e n t and p r e c i p i t a t i o n ha rden ing . )

On the o the r hand, the s t ee l ra i l s m u s t h a v e h i g h s t r eng th coup led w i t h e x c e l l e n t w e a r res i s - tance . M o d e m rai l s t ee l s cons i s t o f a fu l l y pea r l i - t ic m i c r o s t r u c t u r e w i t h a f ine pea r l i t e i n t e r l a m e l - l a r spac ing , as s h o w n in Fig. 2. P e a r l i t e is u n i q u e because i t is a l a m e l l a r c o m p o s i t e c o n s i s t i n g of 88% soft , duc t i l e fe r r i te and 12% hard , b r i t t l e c e m e n t i t e (Fe3C). The ha rd c e m e n t i t e p l a t e s pro- v ide e x c e l l e n t w e a r r e s i s t ance , e s p e c i a l l y w h e n e m b e d d e d in sof t fe r r i te . Pea r l i t i c s t ee l s have h i g h s t r e n g t h and are fu l l y a d e q u a t e to suppo r t h e a v y ax le loads of m o d e m l o c o m o t i v e s and f r e igh t cars. Mos t o f the load is app l i ed in com- p res s ion . Pea r l i t i c s t e e l s a l so have r e l a t i v e l y poor t o u g h n e s s and canno t g e n e r a l l y w i t h s t a n d i m p a c t l o a d s w i t h o u t fa i lu re . The rai l s t ee l cou ld not m e e t the r e q u i r e m e n t s o f the b r idge bu i lde r ,

Table I Mechanical properties of selected steels

Tensile Yield strength strength

Steel Condition MPa ksi MPa kd

Elongation iaS0muma, Reduction Hardness,

Carbon steel bar(a)

1006 Hot rolled 295 43 165 24 30 55 86 Colddrawn 330 48 285 41 20 45 95

1008 Hot rolled 305 44 170 24.5 30 55 86 Colddrawn 340 49 285 41.5 20 45 95

1010 Hot rolled 325 47 180 26 28 50 95 Cold drawn 365 53 305 44 20 40 105

1012 Hot rolled 330 48 185 26.5 28 50 95 Colddrawn 370 54 310 45 19 40 105

1015 Hot rolled 345 50 190 27.5 28 50 101 Cold drawn 385 56 325 47 18 40 111

1016 Hot rolled 380 55 205 30 25 50 110 Cold dmwn 420 61 350 51 18 40 121




1020 Hot rolled 380 55 205 30 25 50 l 1 l Cold drawn 420 61 350 51 15 40 121



1023 Hot rolled 385 56 215 31 25 50 111 Cold drawn 425 62 360 52.5 15 40 121




1527 Hot rolled 515 75 285 41 18 40 149 Colddmwn 570 83 485 70 12 35 163



1536 Hot rolled 570 83 315 45.5 16 40 163 COld drawn 635 92 535 77.5 12 35 187





1541 Hot rolled 635 92 350 51 15 40 187 Cold drawn 705 102.5 600 87 10 30 207 Annealed, cold drawn 650 94 550 80 10 45 184

1042 Hot rolled 550 80 305 44 16 40 163 Colddrawn 6!5 89 515 75 12 35 179 Normalized, cold drawn 585 85 505 73 12 45 179

1043 Hot rolled 565 82 310 45 16 40 163 Cold drawn 625 91 530 77 12 35 179 Normalized, cold drown 600 87 515 75 12 45 179

1044 Hot rolled 550 80 305 44 16 40 163 1045 Hot rolled 565 82 310 45 16 40 163

Colddmwn 625 91 530 77 12 35 179 Annealed, cold drawn 585 85 505 73 12 45 170

1046 Hot rolled 585 85 325 47 15 40 170 Cold drawn 650 94 545 79 12 35 187 Annealed, cold drawn 620 90 515 75 12 45 179

1547 Hot rolled 650 94 360 52 15 30 192 Cold drawn 710 103 605 88 10 28 207 Annealed, cold drawn 655 95 585 85 10 35 187

1548 Hot rolled 660 96 365 53 14 33 197 Colddrawn 735 106.5 615 89.5 10 28 217 Annealed, cold drawn 645 93.5 540 78.5 10 35 192

(continued)

(a) All values are estimated minimum values; type 1100 series steels are rated on the basis of 0.10% max Si or coarse-grain melt- ing practice; the mechanical properties shown are expected minimums for the sizes ranging from 19 to 31.8 mm (0.75 to 1.25 in.). (b) Most data are for 25 mm (1 in.) diam bar. Source: Ref 1

Structure/Property Relationships in Irons and Steels / 155

Table I (continued)


Steel Condition MPa ksi MPa ksi

Elongation in 50 ram, Reduction Hardness,

% ~a area, % HB

Carbon steel bar(a) (continued)

1049 Hot rolled 600 87 330 48 15 35 179 Cold drawn 670 97 560 81.5 10 30 197 Annealed, cold drawn 635 92 530 77 10 40 187

1050 Hot roned 620 90 340 49.5 15 35 179 Cold da'awn 690 100 580 84 10 30 197 Annealed, cold drawn 655 95 550 80 10 40 189

1552 Hot rolled 745 108 410 59.5 12 30 217 Annealed, cold drawn 675 98 570 83 10 40 193

1055 Hot rolled 650 94 355 51.5 12 30 192 Annealed, cold drawn 660 96 560 81 10 40 197

1060 Hot rolled 675 98 370 54 12 30 201 Spheroidized annealed, cold drawn 620 90 485 70 10 45 183

1064 Hot rolled 670 97 370 53.5 12 30 201 Spheroidized annealed, cold drawn 615 89 475 69 10 45 183




1078 Hot rolled 690 1130 380 55 12 30 207 Spheroidized annealed, cold drawn 650 94 500 72.5 10 40 192



1085 Hot rolled 835 121 460 66.5 10 25 248 Spheroidized annealed, cold drawn 695 100.5 540 78 10 40 192

1086 Hot rolled 770 112 425 61.5 10 25 229 Spheroidized aimealed, cold drawn 670 97 510 74 10 40 192






12L14 Hot rolled 395 57 235 34 22 45 121 Cold drawn 540 78 415 60 10 35 163

1108 Hot roUed 345 50 190 27.5 30 50 101 Colddrawn 385 56 325 47 20 40 121


11i7 Hot roned 425 62 235 34 23 47 121 Colddrawn 475 69 400 58 15 40 137


1119 Hot roned 425 62 235 34 23 47 121 Colddrawn 475 69 400 58 15 40 137

1132 Hot roUed 570 83 315 45.5 16 40 167 Cold drawn 635 92 530 77 12 35 183

~1137 Hot roiled 605 88 330 48 15 35 179 Colddrawn 675 98 565 82 10 30 197


1141 Hot roned 650 94 355 51.5 15 35 187 Colddrawn 725 105.1 605 88 10 30 212



1146 Hot roUed 585 85 325 47 15 40 170 Cold drawn 650 94 550 80 12 35 187


(continued)


a n d the H S L A s t r u c t u r a l s tee l c o u l d no t m e e t the r e q u i r e m e n t s o f the c iv i l e n g i n e e r w h o d e s i g n e d the b r i d g e o r the r a i l s y s t e m .

A s i m i l a r c a s e c a n be m a d e fo r the s e l e c t i o n o f c a s t i r o n s . A c a s t m a c h i n e h o u s i n g o n a l a r g e l a t h e r e q u i r e s a m a t e r i a l w i th a d e q u a t e s t r e n g t h , r i g i d i t y , a n d d u r a b i l i t y to s u p p o r t the a p p l i e d l o a d a n d a c e r t a i n d e g r e e o f d a m p i n g c a p a c i t y in o r d e r to r a p i d l y a t t e n u a t e ( d a m p e n ) v i b r a t i o n s f r o m the r o t a t i n g p a r t s o f the l a the . T h e c a s t i r on j a w s o f a c r u s h e r r e q u i r e a m a t e r i a l w i t h s u b s t a n - t ial w e a r r e s i s t a n c e . F o r th i s a p p l i c a t i o n , a c a s t - i n g is r e q u i r e d b e c a u s e w e a r - r e s i s t a n t s t e e l s a r e v e r y d i f f i c u l t to m a c h i n e . F o r the m a c h i n e h o u s - i ng , g r a y ca s t i r on is s e l e c t e d b e c a u s e i t is r e l a - t i v e l y i n e x p e n s i v e , c a n be e a s i l y ca s t , a n d has the a b i l i t y to d a m p e n v i b r a t i o n s as a r e s u l t o f the g r a p h i t e f l a k e s p r e s e n t in i t s m i c r o s t r u c t u r e . T h e s e f l a k e s a r e d i s p e r s e d t h r o u g h o u t the f e r r i t e a n d p e a r l i t e m a t r i x (F ig . 3). T h e g r a p h i t e , b e i n g a m a j o r n o n m e t a l l i c c o n s t i t u e n t in the g r a y i r o n , p r o v i d e s a t o r t u o u s pa th f o r s o u n d to t r a v e l t h r o u g h the m a t e r i a l . W i t h so m a n y f l a k e s , s o u n d w a v e s a re e a s i l y r e f l e c t e d a n d the s o u n d d a m p - e n e d o v e r a r e l a t i v e l y s h o r t d i s t a n c e . H o w e v e r , f o r the j a w c r u s h e r , d a m p i n g c a p a c i t y is no t a r e q u i r e m e n t . In th is c a s e , a n a l l o y w h i t e c a s t i ron is s e l e c t e d b e c a u s e o f i ts h i g h h a r d n e s s a n d w e a r r e s i s t a n c e . T h e w h i t e ca s t i ron m i c r o s t r u c t u r e s h o w n in F ig . 4 is g r a p h i t e f r ee a n d c o n s i s t s o f m a r t e n s i t e in a m a t r i x o f c e m e n t i t e . B o t h o f t hese c o n s t i t u e n t s a re ve ry ha rd a n d thus p r o v i d e the r e q u i r e d w e a r r e s i s t a n c e . T h u s , in t h i s e x a m p l e the g r a y c a s t i r on w o u l d n o t m e e t t he r e q u i r e - m e n t s f o r the j a w s o f a c r u s h e r a n d the w h i t e c a s t i r o n w o u l d n o t m e e t the r e q u i r e m e n t s f o r t h e l a t h e h o u s i n g .

Role of Microstructure In s t ee l s a n d c a s t i r o n s , the m i c r o s t r u c t u r a l

c o n s t i t u e n t s h a v e the n a m e s f e r r i t e , p e a r l i t e , b a i n i t e , m a r t e n s i t e , c e m e n t i t e , a n d a u s t e n i t e . In m o s t a l l o t h e r m e t a l l i c s y s t e m s , the c o n s t i t u e n t s a r e n o t n a m e d , b u t a r e s i m p l y r e f e r r e d to b y a G r e e k l e t t e r (ct, 13, Y, e tc . ) d e r i v e d f r o m the l o c a - t i on o f the c o n s t i t u e n t on a p h a s e d i a g r a m . Fe r - r o u s a l l o y c o n s t i t u e n t s , on the o t h e r h a n d , h a v e b e e n w i d e l y s t u d i e d f o r m o r e t h a n 100 y e a r s . In the e a r l y d a y s , m a n y o f the i n v e s t i g a t o r s w e r e p e t r o g r a p h e r s , m i n i n g e n g i n e e r s , a n d g e o l o g i s t s . B e c a u s e m i n e r a l s h a v e l o n g b e e n n a m e d a f t e r t h e i r d i s c o v e r e r o r p l a c e o f o r i g i n , i t w a s n a t u r a l to s i m i l a r l y n a m e the c o n s t i t u e n t s in s t ee l s a n d c a s t i r ons .

I t c a n b e s e e n t ha t the f o u r e x a m p l e s d e s c r i b e d a b o v e h a v e ve ry d i f f e r e n t m i c r o s t r u c t u r e s : the s t r u c t u r a l s t ee l has a f e r r i t e p lus p e a r l i t e m i c r o - s t r u c t u r e ; the ra i l s t ee l has a f u l l y p e a r l i t i c m i - c r o s t r u c t u r e ; the m a c h i n e h o u s i n g ( l a t he ) has a f e r r i t e p lu s p e a r l i t e m a t r i x w i t h g r a p h i t e f l ake s ; a n d the j a w c r u s h e r m i c r o s t r u c t u r e c o n t a i n s m a r t e n s i t e a n d c e m e n t i t e . In e a c h c a s e , the m i - c r o s t r u c t u r e p l a y s the p r i m a r y r o l e in p r o v i d i n g the p r o p e r t i e s d e s i r e d f o r e a c h a p p l i c a t i o n . F r o m t h e s e e x a m p l e s , o n e c a n see h o w m a t e r i a l p r o p e r - t ies c a n b e t a i l o r e d b y m i c r o s t r u c t u r a l m a n i p u l a - t i o n o r a l t e r a t i o n . K n o w l e d g e a b o u t m i c r o s t r u c - t u r e is t hus p a r a m o u n t in c o m p o n e n t d e s i g n a n d a l l o y d e v e l o p m e n t . In the p a r a g r a p h s t ha t f o l l o w , e a c h m i c r o s t r u c t u r a l c o n s t i t u e n t is d e s c r i b e d w i th p a r t i c u l a r r e f e r e n c e to the p r o p e r t i e s t ha t c a n be d e v e l o p e d b y a p p r o p r i a t e m a n i p u l a t i o n o f the m i c r o s t r u c t u r e t h r o u g h d e f o r m a t i o n (e .g . , ho t a n d c o l d r o l l i n g ) a n d h e a t t r e a t m e n t . F u r t h e r de-

156 / Structure/Property Relationships in Irons and Steels

t a i l s abou t these m i c r o s t r u c t u r a l c o n s t i t u e n t s can be found in R e f 2 to 6.

Ferrite

A w i d e va r i e t y of s t e e l s and cas t i rons fu l l y e x p l o i t the p rope r t i e s o f fer r i te . H o w e v e r , o n l y a f ew c o m m e r c i a l s t ee l s are c o m p l e t e l y fe r r i t i c . An e x a m p l e of the m i c r o s t r u c t u r e of a fu l l y fe r r i t i c , u l t r a l o w ca rbon s t ee l is s h o w n in Fig. 5.

Fe r r i t e is e s s e n t i a l l y a so l id s o l u t i o n of i ron c o n t a i n i n g ca rbon or one or more a l l o y i n g e le- m e n t s such as s i l i con , c h r o m i u m , m a n g a n e s e , and n icke l . There are two types of so l id solu- t ions: i n t e r s t i t i a l and s u b s t i t u t i o n a l . In an in te r - s t i t i a l so l i d so lu t ion , e l e m e n t s w i t h s m a l l a t o m i c d i ame te r , for e x a m p l e , ca rbon and n i t rogen , oc- c u p y spec i f i c i n t e r s t i t i a l s i tes in the b o d y - c e n - t e red cub ic (bcc) i ron c r y s t a l l i n e l a t t i ce . These s i tes are e s s e n t i a l l y the open spaces b e t w e e n the l a r g e r i ron a toms. In a s u b s t i t u t i o n a l so l i d so lu- t ion, e l e m e n t s of s i m i l a r a t o m i c d i a m e t e r r e p l a c e or subs t i t u t e for i ron a toms . The t w o types of so l id so lu t i ons i m p a r t d i f f e r en t c h a r a c t e r i s t i c s to fer r i te . For e x a m p l e , i n t e r s t i t i a l e l e m e n t s l i ke ca rbon and n i t rogen can e a s i l y d i f fuse t h rough the open bcc l a t t i ce , w h e r e a s s u b s t i t u t i o n a l e le- m e n t s l i ke m a n g a n e s e and n i cke l d i f fuse w i t h g rea t d i f f icu l ty . The re fo re , an i n t e r s t i t i a l so l id so lu t ion of i ron and ca rbon r e sponds q u i c k l y during heat t r ea tmen t , w h e r e a s s u b s t i t u t i o n a l so l id s o l u t i o n s b e h a v e s l u g g i s h l y du r ing hea t t reat- men t , such as in h o m o g e n i z a t i o n .

A c c o r d i n g to the i r o n - c a r b o n p h a s e d i a g r a m (Fig . 6a), ve ry l i t t l e c a rbon (0 .022% C) can d is - so lve in fe r r i te (ctFe), e v e n at the e u t e c t o i d t em- pe ra tu re of 727 °C (1330 °F). (The i r o n - c a r b o n phase d i a g r a m i n d i c a t e s the phase r e g i o n s tha t ex i s t ove r a wide ca rbon and t e m p e r a t u r e range . The d i a g r a m r e p r e s e n t s e q u i l i b r i u m cond i t i ons . F i g u r e 6(b) shows an e x p a n d e d i r o n - c a r b o n d ia- g r a m wi th bo th the e u t e e t o i d and e u t e c t i c re- g ions . ) At r o o m t empera tu re , the s o l u b i l i t y is an o rde r of m a g n i t u d e l ess ( b e l o w 0 .005% C). How- ever , even at these s m a l l amoun t s , the a d d i t i o n of ca rbon to pure i ron i n c r e a s e s the r o o m - t e m p e r a - ture y i e l d s t r eng th o f i ron by m o r e than f ive t imes , as seen in Fig. 7. I f the ca rbon c o n t e n t e x c e e d s the s o l u b i l i t y l i m i t o f 0 .022%, the carbon fo rms ano the r p h a s e ca l l ed c e m e n t i t e (Fig . 8). C e m e n t i t e is a l so a c o n s t i t u e n t o f pea r l i t e , as seen in Fig. 9. The ro le o f c e m e n t i t e and p e a r l i t e on the m e c h a n i c a l p r o p e r t i e s of s tee l is d i s c u s s e d be low.

The i n f luence of s o l i d - s o l u t i o n e l e m e n t s on the y i e l d s t r eng th o f fe r r i te is s h o w n in Fig. 10. Here one can c l e a r l y see the s t rong e f fec t o f c a r b o n on i n c r e a s i n g the s t r e n g t h of ferr i te . N i t r o g e n , a l so an i n t e r s t i t i a l e l emen t , has a s i m i l a r effect . Phos- phorus is a lso a fe r r i t e s t r eng thener . In fac t , there are c o m m e r c i a l l y a v a i l a b l e s tee l s c o n t a i n i n g p h o s p h o r u s (up to 0 . 1 2 % P) for s t r e n g t h e n i n g . T h e s e s t ee l s are the r e p h o s p h o r i z e d s t ee l s ( type 1211 to 1215 ser ies) . M e c h a n i c a l p r o p e r t y da ta for these s t ee l s can be found in Table 1.

In Fig. 10, the s u b s t i t u t i o n a l so l id s o l u t i o n e le- m e n t s of s i l i con , copper , m a n g a n e s e , m o l y b d e - num, n i cke l , a l u m i n u m , and c h r o m i u m are s h o w n to have far l e ss e f fec t as fe r r i te s t r e n g t h e n e r s than the in t e r s t i t i a l e l e m e n t s . In fact , c h r o m i u m , n i cke l , and a l u m i n u m in so l id so lu t ion h a v e very l i t t l e i n f l u e n c e on the s t r e n g t h of fer r i te .

In add i t i on to ca rbon (and o ther s o l i d - s o l u t i o n e l emen t s ) , the s t r e n g t h of a fe r r i t i c s tee l is a l so

]'able 1 (continued)

Steel Condition


MPa ksi MPa ksi

Elongatba inSOnma, l~lt~tion Hardm~

% ~a area, % lib

Low-alloy steels(b) 1340 Normalized at 870 °C (1600 °F) 834

Annealed at 800 °C (1475 °F) 703 3140 Normalized at 870 °C (1600 oF) 889

Annealed at 815 °C (1500 °F) 690 4130 Normalized at 870 °C (1600 °F) 670

Annealed at 865 °C (1585 °F) 560 Water quenched from 855 °C (1575 °F) 1040

and tempered at 540 °C (1000 °F) 4140 Normalized at 870 °C (1600 oF) 1020

Annealed at 815 °C (1500 °F) 655 Water quenched from 845 °C ( 1550 °F) 1075

and tempered at 540 °C (1000 °F) 4150 Normalized at 870 °C ( 1600 °F) 1160

Annealed at 830 °C (1525 °F) 731 oil quenched from 830 °C (1525 °F) 1310



Annealed at 810 °C (1490 oF) 745 Oil quenched from 800 °C (1475 °F) 1207



Annealed at 855 °C (1575 oF) 510 4820 Normalized at 860 °C (1580 oF) 758


Annealed at 830 °C (1525 °F) 570 Oil quenched from 845 °C (1550 °F) 972





and tempered at 540 °C (1000 oF) 6150 Normalized at 870 °C (1600 oF) 938


and tempered at 540 °C (1000 oF) 8620 Normalized at 915 °C 0675 °F) 635

Annealed at 870 °C (1600 oF) 540 8630 Normalized at 870 °C (1600 oF) 650

Annealed at 845 °C (1550 °F) 565 Water quenched from 845 °C (1550 °F) 931

and tempered at 540 °C (1000 °F) 8650 Normalized at 870 °C (1600) 1025

Annealed at 795 °C ( 1465 °F) 715 oil quenched from 800 °C (1475 °F) 1185

and tempered at 540 °C ( 1000 °F) 8740 Normalized at 870 °C (1600 oF) 931

Annealed at 815 °C (1500 oF) 696 Oil quenched from 830 °C ( 1525 °F) 1225

and tempered at 540 °C (1000 oF) 9255 Normalized at 900 °C ( 1650 oF) 931


and tempered at 540 °C ( 1000 oF) 9310 Normalized at 890 °C (1630 °F) 910

Annealed at 845 °C (1550 oF) 820

Ferritie stainless steels(b)

405 Annealed bar Cold draw n bar

409 Annealed bar 430 Annealed bar

121 558 81 22.0 63 248 102 434 63 25.5 57 207 129 600 87 19.7 57 262 100 420 61 24.5 51 197 97 435 63 25.5 59.5 197 81 460 67 21.5 59.6 217 151 979 142 18.1 63.9 302

148 655 95 17.7 46.8 302 95 915 60 25.7 56,9 197

156 986 143 15.5 56,9 311

168 731 106 11.7 30,8 321 106 380 55 20.2 40,2 197 190 1215 176 13.5 47.2 375

115 460 67 20.8 51 235 84 425 62 29.0 58 163

186 862 125 12.2 36.3 363 108 470 68 22.0 50.0 217 175 1145 166 14.2 45.9 352

75 350 51 32.5 69.4 143 65 330 48 31.2 62.8 121 83 365 53 29.0 66.7 174 74 370 54 31.3 60.3 149

110 485 70 24.0 59.2 229 99 460 67 22.3 58.8 197

115 470 68 22.7 59.2 229 83 290 42 28.6 57.3 167

141 841 122 18.5 58.9 293

126 530 77 20.7 58.7 255 98 360 52 22.0 43.7 197

159 1000 145 16.4 52.9 311

149 650 94 18.2 50.7 285 105 275 40 17.2 30.6 197 166 1005 146 14.5 45.7 341

136 615 89 21.8 61.0 269 97 415 60 23.0 48.4 197

174 1160 168 14.5 48.2 352

92 360 52 26.3 59.7 183 78 385 56 31.3 62.1 149 94 425 62 23.5 53.5 187 82 370 54 29.0 58.9 156

135 850 123 18.7 59.6 269

149 690 100 14 45.0 302 104 385 56 22.5 46.0 212 172 1105 160 14.5 49.1 352

135 605 88 16.0 47.9 269 101 415 60 22.2 46.4 201 178 1130 164 16.0 53.0 352

135 580 84 19.7 43.4 269 113 485 70 21.7 41.1 229 164 924 134 16.7 38.3 321

132 570 83 18.8 58.1 269HRB 119 450 65 17.3 42.1 241HRB

483 70 276 40 30 60 150 586 85 483 70 20 60 185 450 65 240 35 25 75HRB 517 75 310 45 30 --65" 155

(confnued)

(a) All values are estimated minimum values; type 1100 series steels are rated on the basis of 0.10% max Si or coarse-grain melt- ing practice; the mechanical properties shown are expected minimums for the sizes ranging from 19 to 31.8 mm (0.75 to 1.25 in.). (b) Most data are for 25 mm (1 in.) diam bar. Source: Ref I


Table 1 (continued)

Tensile strength

Steel Ccmdition MPa ksi

Yield strength

MPa ksi

Elongation in 50ram,

%

d e t e r m i n e d b y i ts g r a i n s ize a c c o r d i n g to the H a l l - P e t c h r e l a t i o n s h i p :

Ferritic stainless steels(b) (continued)

430 (cont'd) Annealed and cold drawn 586 85 442 Annealed bar 515 75

Annealed at 815 °C (1500 °F) and cold 545 79 worked

446 Annealed bar 550 80 Annealed at 815 °C (1500 °F) and cold 607 88

drawn

Martensil ic stainless steels(b)

403 Annealed bar 515 75 Tempered bar 765 111

410 Oil quenched from 980 °C ( 1800 °F); 1085 158 tempered at 540 °C (1000 °F);.16 nun (0.625 in.) bar

Oil quenched from 980 °C (1800 °F); 1525 221 tempered at 40 °C (104 °F); 16 mm (0.625 in.) bar

414 Annealed bar 795 115 Cold drawn bar 895 130 Oil quenched from 980 °C (1800 °F); 1005 146

tempered at 650 °C (1200 oF) 420 Annealed bar 655 95

Annealed and cold drawn 760 110 431 Annealed bar 860 125

Annealed and cold drawn 895 130 Oil quenched from 980 °C (1800 °F); 831 121

tempered at 650 °C (1200 oF) Oil quenched from 980 °C (1800 °F); 1435 208

tempered at 40 °C (104 °F) 440C Annealed bar 760 110

Annealed and cold drawn bar 860 125 Hardened and tempered at 315 °C 1970 285

(6OO °F)

Austenitle stainless steels(b)

201 Annealed 760 110 50% hard 1035 150 Full hard 1275 185 Extra hard 1550 225

202 Annealed bar 515 75 Annealed sheet 655 95 50% hard sheet 1030 150

301 Annealed 725 105 50% hard 1035 150 Full hard 1415 205

302 Annealed strip 620 90 25% hard strip 860 125 Annealed bar 585 85

303 Annealed bar 620 90 Colddrawn 690 100

304 Annealed bar 585 85 Annealed and cold drawn 690 100 Cold-drawn high tensile 860 125

305 Annealed sheet 585 85 308 Annealed bar 585 85 309 Annealed bar 655 95 310 Annealed sheet 620 90

Annealed bar 655 95 314 Annealed bar 689 100 316 Annealed sheet 580 84

Annealed bar 550 80 Annealed and cold-drawn bar 620 90

317 Annealed sheet 620 90 Annealed bar 585 85

321 Annealed sheet 620 90 Annealed bar 585 85 Annealed and cold-drawn bar 655 95



(continued)

Reduction Hardness, Gy = Go + kyd -1/2 (Eq 1) in area, % HB

483 70 20 65 185 310 45 30 50 160 427 62 35.5 79 92HRC

345 50 25 45 86HRB 462 67 26 64 96HRB

275 40 35 70 82HRB 585 85 23 67 97HRB

1005 146 13 70 ...

1225 178 15 64 45HRB

620 90 20 60 235 795 115 15 58 270 800 116 19 58 ...

345 50 25 55 195 690 100 14 40 228 655 95 20 55 260 760 110 15 35 270 738 107 20 64 ...

1140 166 17 59 45HRC

450 65 14 25 97HRB 690 100 7 20 260

1900 275 2 10 580

380 55 52 ... 87HRB 760 l l 0 12 ... 32HRC 965 140 8 ... 41HRC

1480 215 1 ... 43HRC 275 40 40 . . . . . . 310 45 40 . . . . . . 760 110 10 _ ... 275 40 60 70' ... 655 95 54 61 ...

1330 193 6 ... 275 40 55 ... 80HRB 515 75 12 _ 25HRC 240 35 60 70" 80HRB 240 35 50 55 160 415 60 40 53 228 235 34 60 70 149 415 60 45 ... 212 655 95 25 ... 275 260 38 50 _ 80HRB 205 30 55 65' 150 275 40 45 65 83HRB 310 45 45 _ 85HRB 275 40 45 65' 160 345 50 45 60 180 290 42 50 _ 79HRB 240 35 60 70" 149 415 60 45 65 190 275 40 45 ... 85HRB 275 40 50 ... 160 240 35 45 _ 80HRB 240 35 55 65' 150 415 60 40 60 185 260 38 40 ... . . . 290 42 45 ... 80HRB 275 40 45 _ 85HRB 240 35 50 65" 160


w h e r e Oy is the y i e l d s t r e n g t h ( in M P a ) , ~o is a c o n s t a n t , ky is a c o n s t a n t , a n d d is the g r a i n d i a m e - te r ( in m m ) .

T h e g r a i n d i a m e t e r is a m e a s u r e m e n t o f s i ze o f the f e r r i t e g r a i n s in the m i c r o s t r u c t u r e , fo r e x a m - p le , n o t e the g r a i n s in the u l t r a l o w c a r b o n s t ee l in F ig . 5. F i g u r e 11 s h o w s the H a l l - P e t c h r e l a t i o n - s h i p f o r a l o w - c a r b o n f u l l y f e r r i t i c s tee l . T h i s r e l a t i o n s h i p is e x t r e m e l y i m p o r t a n t f o r u n d e r - s t a n d i n g s t r u c t u r e - p r o p e r t y r e l a t i o n s h i p s in s t ee l s . C o n t r o l o f g r a i n s i ze t h r o u g h the r - m o m e c h a n i c a l t r e a t m e n t , hea t t r e a t m e n t , a n d / o r m i c r o a l l o y i n g is vi tal to the c o n t r o l o f s t r e n g t h a n d t o u g h n e s s o f m o s t s tee ls . The ro l e o f g r a i n s i ze is d i s c u s s e d in m o r e d e t a i l b e l o w .

T h e r e is a s i m p l e w a y to s t a b i l i z e f e r r i t e , t h e r e b y e x p a n d i n g the r e g i o n o f f e r r i t e in the i r o n - c a r b o n p h a s e d i a g r a m , n a m e l y b y the a d d i - t i o n o f a l l o y i n g e l e m e n t s s u c h as s i l i c o n , c h r o - m i u m , a n d m o l y b d e n u m . T h e s e e l e m e n t s a r e c a l l e d f e r r i t e s t a b i l i z e r s b e c a u s e t h e y s t a b i l i z e f e r r i t e at r o o m t e m p e r a t u r e t h r o u g h r e d u c i n g the a m o u n t o f y so l id s o l u t i o n ( a u s t e n i t e ) w i t h the f o r m a t i o n o f w h a t is c a l l e d a y - l o o p as s e e n at the f a r l e f t in F ig . 12. Th i s i r o n - c h r o m i u m p h a s e d i a - g r a m s h o w s tha t f e r r i t e ex i s t s u p a b o v e 12% C r a n d is s t a b l e up to t he m e l t i n g p o i n t ( l i q u i d u s t e m p e r a t u r e ) . A n i m p o r t a n t f u l l y f e r r i t i c f a m i l y o f s t e e l s is the i r o n - c h r o m i u m f e r r i t i c s t a i n l e s s s t ee l s . T h e s e s tee l s a re r e s i s t a n t to c o r r o s i o n , a n d a r e c l a s s i f i e d as t y p e 4 0 5 , 4 0 9 , 4 2 9 , 4 3 0 , 4 3 4 , 4 3 6 , 4 3 9 , 442 , 444 , a n d 4 4 6 s t a i n l e s s s tee l s . T h e s e s t ee l s r a n g e in c h r o m i u m c o n t e n t f r o m 11 to 3 0 % . A d d i t i o n s o f m o l y b d e n u m , s i l i c o n , n io - b i u m , a l u m i n u m , a n d t i t a n i u m p r o v i d e s p e c i f i c p r o p e r t i e s . F e r r i t i c s t a i n l e s s s t ee l s h a v e g o o d d u c t i l i t y (up to 3 0 % to ta l e l o n g a t i o n a n d 6 0 % r e d u c t i o n in a r e a ) a n d f o r m a b i l i t y , bu t l a c k s t r e n g t h a t e l e v a t e d t e m p e r a t u r e s c o m p a r e d w i t h a u s t e n i t i c s t a i n l e s s s t ee l s . R o o m - t e m p e r a t u r e y i e l d s t r e n g t h s r a n g e f r o m 170 to a b o u t 4 4 0 M P a (25 to 6 4 ks i ) , a n d r o o m - t e m p e r a t u r e t en s i l e s t r e n g t h s r a n g e f r o m 3 8 0 to a b o u t 5 5 0 M P a (55 to 8 0 ks i ) . T a b l e 1 l i s ts t he m e c h a n i c a l p r o p e r t i e s o f s o m e o f t h e f e r r i t i c s t a i n l e s s s t ee l s . T y p e 4 0 9 s t a i n l e s s s tee l is w i d e l y u s e d f o r a u t o m o t i v e ex- h a u s t s y s t e m s . T y p e 4 3 0 f r e e - m a c h i n i n g s t a i n l e s s s t ee l h a s the b e s t m a c h i n a b i l i t y o f a l l s t a i n l e s s s t ee l s o t h e r t h a n t ha t o f a l o w - c a r b o n , f r e e - m a - c h i n i n g m a r t e n s i t i c s t a i n l e s s s tee l ( t y p e 41.6).

A n o t h e r f a m i l y o f s t ee l s u t i l i z i n g a f e r r i t e s t a - b i l i z e r ( y - l o o p ) a r e t he i r o n - s i l i c o n f e r r i t i c a l l o y s c o n t a i n i n g u p to a b o u t 6 . 5 % Si ( c a r b o n - f r e e ) . T h e s e s t ee l s a r e o f c o m m e r c i a l i m p o r t a n c e be - c a u s e t h e y h a v e e x c e l l e n t m a g n e t i c p e r m e a b i l i t y a n d l o w c o r e loss . H i g h - e f f i c i e n c y m o t o r s a n d t r a n s f o r m e r s a re p r o d u c e d f r o m t h e s e i r o n - s i l i - c o n e l e c t r i c a l s t e e l s ( a l u m i n u m c a n a l s o subs t i - t u t e f o r s i l i con in t hem) .

O v e r t he p a s t 20 y e a r s o r so , a n e w b r e e d o f v e r y - l o w - c a r b o n f u l l y f e r r i t i c s h e e t s t ee l s h a s e m e r g e d fo r a p p l i c a t i o n s r e q u i r i n g e x c e p t i o n a l f o r m a b i l i t y ( see F ig . 5). T h e s e a re the i n t e r s t i - t i a l - f r e e ( IF) s t ee l s fo r w h i c h c a r b o n a n d n i t ro - g e n a re r e d u c e d in the s t e e l m a k i n g p r o c e s s to v e r y l o w l eve l s , a n d a n y r e m a i n i n g i n t e r s t i t i a l c a r b o n or n i t r o g e n is t i ed u p w i t h s m a l l a m o u n t s o f a l l o y i n g e l e m e n t s (e .g . , t i t a n i u m o r n i o b i u m ) t h a t f o r m p r e f e r e n t i a l l y c a r b i d e s a n d n i t r i d e s .


Table I (continued)

qI~mBe Yield

st~ngth strength Sted Cand~laa MPa k~ MPa I~

Elongation inS0mm, Reduction Hardness,

% in area, % liB

Austenilic stainless steels(b) (continued)

347 (eont'd) Annealedandcolddrawnbar 690 384 Annealed wire 1040 °C (1900 °F) 515

Maraging steels(b)

18Ni(250) Annealed 965 Aged bar 32 mm (1.25 in.) 1844 Aged sheet 6 mm (0.25 in.) 1874

18Ni(300) Annealed 1034 Aged bar 32 mm (1.25 in.) 2041 Aged sheet 6 mm (0.25 in.) 2169

18Ni(350) Annealed 1140 Aged bar 32 mm (l.25 in.) 2391 Aged sheet 6 mm (0.25 in.) 2451

100 450 65 40 60 212 75 240 35 55 72 70HRB

140 655 95 17 75 30 HRC 269 1784 259 11 56.5 51.8 HRC 272 1832 266 8 40.8 50.6HRC 150 758 110 18 72 32HRC 296 2020 293 11.6 55.8 54.7 HRC 315 2135 310 7.7 35 55.1HRC 165 827 120 18 70 35 HRC 347 2348 341 7.6 33.8 58.4 HRC 356 2395 347 3 15.4 57.7 HRC

(a) All values are estimated minimum values; type 1100 series steels ate rated on the basis of 0.10% max Si or coarse-grain melt- ing practice; the mechanical properties shown are expected minimums for the sizes ranging from 19 to 31.8 mm (0.75 to 1.25 in.). (b) Most data are for 25 mm (1 in.) diam bar. Some: Ref 1

These steels have very low strength, but are used to produce components that are diff icul t or im- possible to form from other steels. Very-low-carbon, fully ferritic s teels (0.001% C) are now being manufac tured for automot ive components that harden during the paint -cur ing cycle. These s teels are called bake-hardening steels and have controlled amounts o f carbon and ni t rogen that combine with other e lements , such as t i tan ium and niobium, during the baking cycle (175 °C, or 350 °F, for 30 min). The process is called aging, and the s trength derives f rom the precipitat ion o f t i t an ium/niobium carbonitr ides at the elevated temperature.

Another form of very- low-carbon, fully ferritic steel is motor laminat ion steel. The carbon is re- moved from these steels by a process known as decarburization. The decarburized (carbon-free) ferritic steel has good permeabi l i ty and suffi- ciently low core loss (not as low as the iron-sil i- con alloys) to be used for electric motor lamina-

tions, that is, the stacked steel layers in the rotor and stator o f the motor.

As noted previously, a number of properties are exploited in fully ferritic steels:

• I ron-s i l i con s teels: Exceptional electrical propert ies

• I r o n - c h r o m i u m steels: Good corrosion resistance

• In ters t i t ia l - f ree s teels: Exceptional formability

• B a k e - h a r d e n i n g s teels: Strengthens during paint cure cycle

• Lamina t ion s tee ls : Good electrical properties

PearlRe

As the carbon content of steel is increased be- yond the solubili ty l imit (0.02% C) on the iron- carbon binary phase diagram, a const i tuent called

pearli te forms. Pearlite is formed by cooling the steel through the eutectoid temperature (the temperature o f 727 °C in Fig. 6) by the following reaction:

Austenite ~ cementite + ferrite ffXl2)

The cement i te and ferrite form as parallel plates called lamellae (Fig. 13). This is essent ial ly a compos i t e mic ros t ruc tu re cons i s t ing o f a very hard carbide phase, cementi te, and a very soft and ductile ferrite phase. A fully pearlitic microstructure is formed at the eutectoid composi t ion of 0.78% C. As can be seen in Fig. 2 and 13, pearlite forms as colonies where the lamellae are al igned in the same orientation. The propert ies of fully pearlitic steels are determined by the spacing between the ferr i te-cementi te lamellae, a d imension called the inter lamellar spacing, X, and the colony size. A simple relat ionship for yield s trength has been developed by Heller (Ref 10) as follows:

fly = -85.9 + 8.3 (X -t/2) (Eq 3)

where fly is the 0.2% offset yield s trength (in MPa) and X is the interlamellar spacing (in mm). Figure 14 shows Hel ler ' s plot of s t rength versus inter lamellar spacing for fully pearlitic eutectoid steels.

It has also been shown by Hyzak and Bernstein (Ref 11) that s t rength is related to inter lamellar spacing, pearlite colony size, and prior-austenite grain size, according to t h e following relationship:

YS = 52.3 + 2.18 (~-1/2) - 0 . 4 (de -L'2) -2 .88 (d-1/2)(Eq 4)

where YS is the yield s trength (in MPa), d e is the pearli te colony size (in mm), and d is the prior- austeni te grain size (in mm). From Eq 3 and 4, it can be seen that the steel composi t ion does not have a major inf luence on the yield strength of a fully pearlitic eutectoid steel. There is some solid-

Fig, 3 Microstructure of a gray cast iron with a ferrite-pearlite matrix. Note the graphite Fig. 4 Microstructure of an alloy white cast iron. White constituent is cementite and the flakes dispersed throughout the matrix. 4% picral etch. 320x. Courtesy of A.O. darker constituent is martensite with some retained austenite. 4% picral etch.

Benscoter, Lehigh University 250x. Courtesy ofA.O. Benscoter, Lehigh University

Structure/Property Relationships in Irons and S tee ls /159

Fig. 5 Microstructure of a fully ferritic, ultralow carbon steel. Marshalls etch + HF, 300x. Courtesy of

A.O. Benscoter, Lehigh University

solution s t rengthening of the ferrite in the lamellar structure (see Fig. 10).

The thickness of the cement i te lamellae can also influence the propert ies of pearlite. Fine cementi te lamellae can be deformed, compared with coarse lamellae, which tend to crack during deformation.

Al though fully pearlitic steels have high strength, high hardness , and good wear resistance, they also have poor ductil i ty and toughness. For example, a low-carbon, fully ferritic

steel will typically have a total elongation of more than 50%, whereas a fully pearlitic steel (e.g., type 1080) will typically have a total elongation of about 10% (see Table 1). A low-carbon fully ferritic steel will have a room-tempera ture Charpy V-notch impact energy of about 200 J (150 f t . lbf), whereas a fully pearlitic steel will have room-tempera ture impact energy of under 10 J (7 f t . lbf). The transit ion temperature (i.e., the temperature at which a material changes f rom ductile fracture to brittle fracture) for a ful ly pearlitic steel can be approximated from the following relat ionship (Ref 11):

TT = 217.84 - 0.83 (de -1/2) - 2.98(d -1"~) (Eq5)

where TT is the transit ion temperature (in °C). From Eq 5, one can see that both the prior-

austeni te grain size and pearlite colony size control the transi t ion temperature of a pearlitic steel. Unfortunately, the transit ion temperature of a fully pearlitic steel is always well above room temperature. This means that at room temperature the general fracture mode is cleavage, which is associated with brittle fracture. Therefore, fully pearlitic steels should not be used in applicat ions where toughness is important. Also, pearlitic s teels with carbon contents sl ightly or mod- erately higher than the eutectoid c o m p o s i t i o n (called hypereutectoid steels) have even poorer toughness .

From Eq 4 and 5, one can see that for pearlite, s trength is controlled by interlamellar spacing, colony size, and prior-austenite grain size, and toughness is controlled by colony size and prior-

austeni te grain size. Unfortunately, these three factors are rather difficult to measure . To determine inter lamellar spacing, a scanning electron microscope (SEM), or a t ransmiss ion electron microscope (TEM) is needed in order to resolve the spacing, Generally, a magnif ica t ion of 10,000x is adequate, as seen in Fig. 13. Special stat ist ical procedures have been developed to determine an accurate measu remen t o f the spacing (Ref 12). The colony size and especial ly the pr ior-austeni te grain size are very difficult to measu re and require a skilled meta l lographer us- ing the l ight microscope or SEM and special e tching procedures.

Because of poor duct i l i ty / toughness , there a r e

only a few applicat ions for fully pearlitic steels, inc luding railroad rails and wheels and high- s t rength wire. By far, the largest tonnage applica- t ion is for rails. A fully pearlitic rail steel pro- vides excel lent wear resis tance for r a i l road wheel/rai l contact. Rail life is measured in mil- l ions of gross tons (MGT) of travel and current rail life easi ly exceeds 250 MGT. The wear resistance of pearlite arises f rom the unique morphology of the ferr i te-cementi te lamel lar composi te where a hard const i tuent is embedded into a soft- ducti le const i tuent . This means that the hard cement i te plates do not abrade away as easily as the rounded cement i te part icles found in other steel micros t ruc tures , that is, tempered mar tens i te and bainite, which is d iscussed later. Wear resis tance o f a rail steel is directly proportional to hardness . This is shown in Fig. 15, which indicates less weight loss as hardness increases. Also, wear re- s is tance (less weight loss) increases as inter- lamel lar spacing decreases, as shown in Fig. 16.

¢D O. E

1180

1140

1100

1060

1020

980

940

900

86O

820

780

740

700 /

66O Fe

Carbon, at.%

1 2 3 4 5 6 I I I I I I

Fe-C equilibrium (experimental)

- - Fe-Fe3C equilibrium (experimental)

(~Fe) auatenite

~ 9 1 2 °C , / "

~ ~ 0 ¢ F . ) ferrite . , - ~

% ~ 770 °C (Curie temperature) -*°~ . . . / . . . . . . . . . . . . . . . . . . ~- - - ' - ~ 0.68 7 I ~ 0.0206 ~ ~ , . ' ° "

0.0218 I I Ferrite + cementite

I I I I 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 1.4

7 8 9

1154°C - ~ . . . ~ 2125

I J . , ~1 , , / 2.08 ~ "'" 8 °C-'~ 2050

. o " Y 211 -- 1975

• *' Y • . ~ -- 1900

-- 1825

-- 1750

AUS tenite + cementite -- 1700

-- 1625

-- 1550

o u-

E

-- 1475

738 °C - 1400

I - - 1325 727 °C

I - 1250

1.5 1.6 1.7 1.8 1.9 2.0 2.1 2.2

Carbon, wt%

Fig. 6 (a ) Iron-carbon phase diagram showing the austenite (y Fe) and ferrite (ocFe) phase regions and eutectoid composition and temperature. Dotted lines represent iron-graphite equilibrium conditions and solid lines represent iron-cementite equilibrium conditions. Only the solid lines are important with respect to steels. Source: Ref 2


Thus, the mos t important micros t ructura l parameter for control l ing hardness and wear resistance is the pearlite inter lamellar spacing. Fortu- nately, inter lamellar spacing is easy to control and is dependent solely on t ransformat ion temperature.

Figure 17 shows a cont inuous cooling transfor- mat ion (CCT) d iagram for a typical rail steel. A CCT diagram is a t ime versus temperature plot showing the regions at which various consti tu- c n t s - - f e r d t e , pearlite, bainite, and m a r t e n s i t e - - form during the cont inuous cooling of a steel component . Usual ly several cool ing curves are shown with the associa ted start and f inish transformation temperatures of each const i tuent . These d iagrams should not be confused with isothermal t ransformat ion (IT or TTT) d iagrams, which are derived by rapidly quenching very thin spec imens to various temperatures , and mainta in- ing that temperature ( isothermal) until the speci- mens begin to t ransform, partially t ransform, and fully t ransform, at which t ime they are quenched to room temperature. An IT d iagram does not represent the t ransformat ion behavior in mos t

processes where steel parts are cont inuously cooled, that is, air cooled, and so forth.

As shown in Fig. 17, the pead i t e t ransformation temperature (indicated by the pearli te-start curve, Ps) decreases with increasing cooling rate. The hardness of peaflite increases with decreas- ing t ransformat ion temperature. Thus , in order to provide a rail steel with the h ighes t hardness and wear resis tance, one mus t cool the rail f rom the austeni te at the fastest rate possible to obtain the lowest t ransformat ion temperature. This is done in practice by a process known as head hardening, which is s imply an accelerated cooling process us ing forced air or water sprays to achieve the desired cooling rate (Ref 15). Because only the head of the rail contacts the wheel of the railway car and locomotive, only the head re- quires the higher hardness and wear resistance.

Another application for a fully pearlitic steel is h igh-s t rength wire (e.g., piano wire). Again, the composi te morphology of lamellar ferrite and cement i te is exploited, this t ime during wire drawing. A fully pearlitic steel rod is heat treated by a process known as patenting. Dur ing patenting,

1~ M 3270

1~ 3090

2910

1 ! 2730

GFe 1~ 2550

2370

2190

11 2010

lC 1830 ~

E i ~ E 1470

7 1290

rILE[

-~ 930

4 750

3 570

I P_~

30 Fe 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0

Carbon, wt%

Fig. 6(b) Expanded iron-carbon phase diagram showing both the eutectoid (shown in Fig. 6a) and eutectic regions. Dotted lines represent iron-graphite equilibrium conditions and solid lines represent iron-cementite equilib-

rium conditions. The solid lines at the eutectic are important to white cast irons and the dotted lines are important to gray cast irons. Source: Ref 2

the rod is t ransformed at a temperature of about 540 °C (1000 °F) by passing it through a lead or salt bath at this temperature. This develops a micros t ructure with a very fine pearli te inter- lamel lar spacing because the t ransformation takes place at the nose of the CCT diagram, that is, at the lowest possible pearlite t ransformation temperature (see Fig. 17). The rod is then cold drawn to wire. Because o f the very fine inter- lamel lar spacing, the ferrite and cementi te lamellae become al igned along the wire axis during the deformat ion process. Also, the fine ccmenti te lamel la tend to bend and deform as the wire is e longated during drawing. The resul t ing wire is one of the s t rongest commercia l products avail- able; for example , a commercia l 0.1 m m (0.004 in.) diam wire can have a tensile s t rength in the range of 3.0 to 3.3 GPa (439 to 485 ksi), and in special cases a tensile s t rength as h igh as 4.8 GPa (696 ksi) can be obtained. These wires are used in musica l ins t ruments because of the sound quality developed from the high tensi le stresses applied in s tr inging a piano and violin and are also used in wire rope cables for suspension bridges.

Ferrite-Pearlite

The mos t common structural s teels produced have a mixed ferrite-pearli te microstructure. Their applicat ions include beams for bridges and high-r ise buildings, plates for ships, and reinforcing bars for roadways. These steels are rela- t ively inexpensive and are produced in large ton- nages. They also have the advantage o f being able to be produced with a wide range of properties. The micros t ructure of typical ferrite-pearlite steels is shown in Fig. 18.

In mos t ferrite-pearlite steels, the carbon content and the grain size determine the microstructure and result ing properties. For example, Fig. 19 shows the effect of carbon on tensi le and impact properties. The ult imate tensile strength steadily increases with increasing carbon content. This is caused by the increase in the volume fraction o f pearlite in the microstructure , which has a s t rength much higher than that of ferrite. Thus , increas ing the volume fraction o f pearlite has a profound effect on increasing tensile strength.

However, as seen in Fig. 19, the yield strength is relatively unaffected by carbon content, r ising f rom about 275 MPa (40 ksi) to about 415 MPa (60 ksi) over the range of carbon content shown. This is because yielding in a ferrite-pearlite steel is controlled by the fcrrite matrix, which is generally considered to be the cont inuous phase (ma-

O3

"~ 35 241 ~:

-'~ 25 ~ ' ..... 172

"N, / 103= 0

o~ 10 ~ o. 0 0.001 0.002 0.003 0.004 0.005 o o 6

Carbon, wt%

Fig, 7 Increase in room-temperature yield strength of iron with small additions of carbon. Source: Ref 7


Fig. 8 Photomic.rograph of an annealed low-carbon sheet steel with grain-boundary cementite. 2% nital + 4% picral etch. 1000x

Fig. 9 Photomicrograph of pearlite (dark constituent) in a low-carbon steel sheet. 2% nital + 4% picral etch. 1000x

trix) in the microstructure. Therefore, pearlite plays only a minor role in yielding behavior.

From Fig. 19, one can also see that ductility, as represented by reduct ion in area, steadily decreases with increas ing carbon content. A steel with 0.10% C has a reduction in area of about 75%, whereas a steel with 0.70% C has a reduction in area of only 25%. Percent total elongation would show a s imilar trend, however, with values much less than percent reduction in area.

Much work has been done to develop empirical equations for ferri te-pearli te steels that relate strength and toughness to microstructural features, for example , grain size and percent of pearlite as well as composi t ion. One such equation for ferri te-pearli tc steels under 0.25% C is as fol lows (Ref 16):

YS = 53.9 + 32.34 (Mn) + 83.2(Si) + 354.2(Nf) + 17.4(d-U2) (Eq 6)

where Mn is the manganese content (%), Si is the si l icon content (%), Nf is the free ni t rogen content (%), and d is the ferrite grain size (in mm). Equa- tion 6 shows that carbon content (percent pearlite)

has no effect on yield strength, whereas the yield s t rength in Fig. 19 increases somewhat with carbon content. According to Eq 6, manganese , silicon, and ni t rogen have a pronounced effect on yield strength, as does grain size. However, in most ferri te-pearli te steels ni trogen is quite low (under 0.010%) and thus has minimal effect on yield strength. In addition, as discussed below, ni t rogen has a detrimental effect on impact properties.

The regress ion equation for tensile s t rength for the same steels is as fol lows (Ref 16):

TS = 294,1 + 27.7(Mn) + 83.2(Si) + 3.9(P) + 7.7(d -lt2) (F-4 7)

where TS is the tensile s t rength (in MPa) and P is pearli te content (%). Thus , in dist inction to yield s trength, the percentage o f pearlite in the micro- s t r u c t u r e has an i m p o r t a n t e f f ec t on t e n s i l e s trength.

Toughness of ferrite-pearlite steels is also an important considerat ion in their use. It has long been known that the absorbed energy in a Charpy V-notch test is decreased by increasing carbon content, as seen in Fig. 20. In this graph of im-

pact energy versus test temperature, the shel f energy decreases f rom about 200 J (150 ft • lbf) for a 0.11% C steel to about 35 J (25 f t . lbf) for a 0.80% C steel. Also, the transi t ion temperature increases f rom about - 5 0 to 150 °C ( -60 to 300 °F) over this same range o f carbon content. The effect of carbon is due mainly to its effect on the percentage of pearlite in the microstructurc . This is reflected in the regress ion equat ion for transi- t ion temperature below (Ref 16):

TT = -19 + 44(Si) + 700(N~/2)

+ 2.2(P) - 11.5 (d -1/2) (F_.q 8)

It can be seen in all these re la t ionships that ferrite grain size is an important parameter in improving both s trength and toughness . It can also be seen that while pearlite is beneficial for increasing tensi le s t rength and n i t rogen is beneficial for increasing yield strength, both are harm- ful to toughness . Therefore, me thods to control the grain size of ferri te-pearli te steels have rapidly evolved over the past 25 years. The two mos t important methods to control grain size are controlled roll ing and microal loying. In fact, these

4-375

+225

.--~_m+150

"~ +75

o 0

-75

I C and N

Si

/ y - - Ni and AI

0 0.5 1.0 1.5 2.0 2.5 3.0 Alloy content, wt%

Fig, 1 0 Influence of solid-solution elements on the changes in yield stress of low-carbon ferritic

steels. Source: Ref 5

600

500

400 &

~ 300

200

100

Fig. 11

I I I I I I I I I I I I 0 1 2 3 4 5 6 7 8 9 10 11 12

Grain diameter (d-l~), mm -1~

Hall-Petch relationship in low-carbon ~mtic steels, souse: Ref 8

80

80

"N. 20 |


oo

(9 ¢:L E

Fig. 12

Chromium, at.%

0 10 20 30 40 50 60 70 20OO

1800

1600

I I

80 90 100

1538 °C ~ . . . . . .

21 1400 - 1394 °C

1200 - ~ (~Fe,Cr)

1000 _ ( ~ F e ) / / _ 1 2 . 7 oc/I

oc 8001:-- -.7

I - / ( o I ~nn I Magnetic " ~ • "---- . . . . . I , "* " .

Itransformabon.- ,, : " - . . 475 o C

o.'. . . . . . . . . . . . . . . . . . . =.." . . . . . . . . . . . . . . . . . . . . . . . . . . . ".,.. / 400 i r'1 °° I I I I I I I t "'~

0 10 20 30 40 50 60 70 80 90

Fe Chromium, wt%

I i I I I I I

1863 °C

1516 ° :

100

Cr

Iron-chromium phase diagram. Source: Ref 9

methods are used in conjunction to produce strong, tough ferrite-pearlite steels.

Controlled rolling is a thermomechanical treatment in which steel plates are rolled below the recrystailization temperature of aastcnite. This process results in elongation of the austenite grains. Upon further rolling and subsequent cooling to room temperature, the austenite-to-ferrite transformation takes place. The ferrite grains are restricted in their growth because of the "pancake" austeaite grain morphology. This produces the fine ferrite grain size required for higher strength and toughness.

Microalloying is the term applied to the addition of small amounts of special alloying elements (vanadium, niobium, or titanium) that aid

in retarding austenite recrystallization, thus al- lowing a wide window of rolling temperatures for controlled rolling. Without retarding recrystallization, as in normal hot rolling, the pancake- type grains do not form and a fine grain size cannot be developed. Microalloyed steels are used in a wide variety of high tonnage applications including structural steels for the construc- tion industry (bridges, multistory buildings, etc.), reinforcing bar, pipe for gas transmission, and numerous forging applications.

Bainite

Like pearlite, bainitc is a composite of ferrite and cementitc. Unlike pearlite, the ferritc has an

Fig, 13 SEM micrograph of pearlite showing ferrile and cementite lamellae. 4% picral etch. 10, O00x

acicular morphology and the carbides are discrete particles. Because of these morphological differences, bainite has much different property characteristics than pearlite. In general, bainitic steels have high strength coupled with good toughness, whereas pearlitic steels have high strength with poor toughness.

Another difference between baiaite and pearlite is the complexity of the bainite morphologies compared with the simple lamellar morphology of pearlite. The morphologies of bainite are still being debated in the literature. For years, since the classic work of Bain and Davenport in the 1930s (Ref 18), there were two classifications of bainite: upper and lower bainite. This nomencla- ture was derived from the temperature regions at which bainite formed during isothermal (constant temperature) transformation. Upper bainite formed isothermally in the temperature range of 400 to 550 °C (750 to 1020 °F), and lower bainite formed isothermally in the temperature range of 250 to 400 °C (480 to 750 °F). Exam- ples of the microstructure of upper and lower bainite are shown in Fig. 21. One can see that both types of bainite have an acicular morphology, with upper bainite being coarser than lower bainite. The true morphological differences between the microstructures can only be determined by electron microscopy. Transmission electron micrographs of upper and lower baiaite are shown in Fig. 22. In upper bainitc, the iron carbide phase forms at the lath boundaries, whereas in lower bainite, the carbide phase forms on particular crystallographic habit planes within the laths. Because of these differences in morphology, upper and lower bainite have different mechanical properties. Lower bainite, with a fine acicular structure and carbides within the laths, has higher strength and higher toughness than upper bainite with its coarser structure.

Because during manufacture most steels un- dergo continuous cooling rather than isothermal holding, the terms upper and lower baiaite can become confusing because "upper" and "lower" are no longer an adequate description of morphology. Bainite has recently been reclassified by its morphology, not by the temperature range in which it forms (Ref 19). For example, a recent classification of bainite yields three distinct types of morphology.

Class 1 (B1): Acicular ferrite associated with intralath (plate) iron carbide, that is, cemcn- tite (replaces the term "lower bainite")

900 O .

8O0

~ 7oo

.~>6OO

~.500 0

Interlamellar spacing (Sp), nm

300 200 100 80 60 I I f I

S . j ¢ ,-

400 60 80 100 120 140

Reciprocal root of Interlamellar spacing (Sp-1/2), mm -1/2

Fig. 14 Relationship behveen peadite interlamellar spacing and yield strength for eutectoid steels.

Source: Ref I0

Structure/Property Relationships in Irons and Steels/163

o ~

o J ~

o

Fig. 15

1.6

1.2

0.8

0.4

0 2OO 225 250 275 300 325 350 375

Brinell hardness, HB

Relationship between hardness and wear resistance (weight loss) for rail steels. Source: Ref 13

1.6

/ 1,2

o= j z 0.8 / ~

i < 0.4

0 0.06 0.08 0.10 0.12 0.14 0.16 0.18 0.20 0.22 0.24 0.26

Fig. 16

Pearlite spacing, pm

Relationship between pearlite interlamellar spacing and wear resistance (weight loss) for rail steels. Source: Ref 13

• Class 2 (B2): Acicular ferrite associated with in ter lath (plate) part icles or f i lms of cement i te and/or austeni te (replaces the term "upper bainite")

• Class 3 (B3): Acicular ferrite associa ted with a const i tuent consis t ing of discrete is lands of austenite and/or martensi te

The bainitic steels have a wide range of mechanical properties depending on the microstructural morphology and composi t ion; for example, yield strength can range f rom 450 to 950 MPa (65 to 140 ksi), and tensile s t rength from 530 to 1200 MPa (75 to 175 ksi). Another aspect of a bainitic steel is that a single composit ion, I/2Mo-B steel for example, can yield a bainitic microstructure over a wide range o f t ransformation temperatures. The CCT diagram for this steel is shown in Fig. 23. Note that for this steel the bainite start (Bs) temperature is a lmost constant at 600 °C (1110 °F). This flat t rans formation region is important because t ransformation temperature plays an important role in the development of microstructure . A constant t ransformation temperature permits the development of a similar microstructure and properties over a wide range of cooling rates. This has many advantages in the manufac tur ing of bainitic steels and is particularly advantageous in thick sect ions where a wide range in cooling rates is found f rom the surface to the center of the part.

In designing a bainitic steel with a wide transformation region, it becomes crit ical that the pearlite and ferrite regions are pushed as far to the right as possible on the CCT diagram; that is, pearlite and ferrite form only at slow cooling rates. Alloying e lements such as nickel, chromium, and mo lybdenum (and manganese) are selected for this purpose.

For low-carbon bainitic steels, the relat ionship between t ransformat ion temperature and tensi le strength is shown in Fig. 24 (martensi te is discussed in the next section). Note the rapid increase in tensi le s t rength as the t ransformation temperature decreases. For these steels, a regres- sion equation for tensi le s t rength has been developed as follows (Ref 21):

TS = 246.4 + 1925(C) + 231(Mn + Cr) + 185(Mo) + 92(W) + 123(Ni) + 62(Cu) + 385(V + 11) (Eq9)

In addition to the e lements carbon, nickel, chromium, molybdenum, vanadium, and so forth, it is well known that boron in very small quanti-

ties (for example, 0.003%) has a pronounced effect on retarding the ferrite t ransformation. Thus , in a boron-containing steel (e.g., l/2Mo + B), the ferri te nose in the CCT diagram is pushed to s lower cooling rates. Boron retards the nuclea- tion of ferrite on the austenite grain boundar ies and, in doing so, permits bainite to be formed (Fig. 23). Whenever boron is added to steel, it mus t be prevented from combining with other e lements such as oxygen and nitrogen. Generally, a l umi num and t i tanium are added first in order to lower the oxygen and nitrogen levels o f the steel. Even when adequately protected, the effective- ness of boron decreases with increasing carbon content and austeni te grain size.

At tempts have been made to quant i ta t ively relate the microstructural features o f bainite to mechanical properties. One such relat ionship is (Ref 22):

YS = -194 + 17.4(d -1/2) + 15(nl/4) (Eq 10)

where YS is the 0.2% offset yield s trength (in MPa), d is the bainite lath size (mean l inear inter-

cept) (in ram), and n is the nu m b er o f carbides per m m 2 in the plane of section.

With bainit ic steels, the lath width of the bainite obeys a Hall-Petch re la t ionship as shown in Fig. 25. The lath size is directly related to the austeni te grain size and decreases with decreas- ing bainite t ransformat ion temperature. Because of the fine microstructure of bainite, the meas- u rement of lath size and carbide densi ty can only be done by SEM or TEM.

In low-carbon bainitic steels, type B 2 (upper) bainite has inferior toughness to type B 1 (lower) bainite. In both cases, s t rength increases as the t ransi t ion temperature decreases. In type B 2 (upper) bainite, the carbides are m u c h coarser than in type B 1 (lower) bainite and have a tendency to crack and initiate c leavage (brittle) fracture. In type B l bainite, the small carbides have less tendency to fracture. One can lower the transi t ion temperature in type B l bainitic s teels by provid- ing a finer austenite grain size through lower- tempera ture thermomechanica l t reatment and grain ref inement .

Bainit ic steels are used in m an y applicat ions inc luding pressure vessels , backup rolls, turbine

9oo

400

300

200

100

0 10

700

? 600

¢ 500 t z E

_I o.n,r .e °Clmln ~ ~s

~_ 1 -643 \ \ 1 ~ ~ 2 - 600 0-545_ ( % ~ . 'Bf 4 - 500 5 - 450 6 400 7 - 352 M s \ "~ 8 - 300 9 - 253

10 - 225 Martensite 11 -189 12 - 50

I I I I I I

lO0 Time, s

1000

Fig. 17 ^ ccT diagram of a typical rail steel (composition: 0.77% C, 0.95% Mn, 0.22% Si, 0.014% P, 0.017% S, 0.10% Cr). Source: Ref 14


(a) (h)

Fig. 18 Microstructure of typical ferrite-pearlite structural steels at two different carbon contents. (a) 0.10% C. (b) 0.25% C. 2% nital + 4% picral etch. 200x

rotors, die blocks, d ie-cast ing molds, nuclear re- actor components , and ear thmoving equipment . One major advantage of a bainitic steel is that an optimal s t rength / toughness combinat ion can be produced without expens ive heat t reatment , for example , quenching and tempering as in martensitic steels.

Martensite

Martensi te is essent ial ly a supersaturated solid solution o f carbon in iron. The amount of carbon in mar tens i te far exceeds that found in solid solution in ferrite. Because of this, the normal body- centered cubic (bcc) lattice is distorted in order to accommodate the carbon atoms. The distorted lattice becomes body-centered tetragonal (bet). In pla in-carbon and low-alloy steels, this super- saturat ion is general ly produced through very rapid cooling f rom the austeni te phase region (quenching in water, iced-water, brine, iced- brine, oil or aqueous po lymer solut ions) to avoid forming ferrite, pearlite, and bainite. Some highly al loyed steels can form martensi te upon air cooling (see the d iscuss ion o f marag ing steels later in this section). Depending on carbon content, martensi te in its quenched state can be very hard and brittle, and, because of this bri t t leness, martensi t ic steels are usual ly tempered to restore some ducti l i ty and increase toughness .

Reference to a CCT diagram shows that mar tens i te only forms at high cooling rates in pla in-carbon and low-al loy steels. A CCT dia- g ram for type 4340 is shown in Fig. 26, which indicates that mar tens i te forms at cooling rates exceeding about 1000 °C/rain. Most commercia l martensi t ic steels contain deliberate al loying addit ions intended to suppress the format ion o f other cons t i t uen t s - - tha t is ferrite, peariite, and ba in i t e - -du r ing cont inuous cooling. This means that these const i tuents form at s lower cooling rates, a l lowing mar tens i te to form at the faster cooling rates, for example , during oil and water quenching. This concept is called hardenabil i ty and is essent ial ly the capaci ty o f a steel to harden

by rapid quenching. Most all the convent ional al loying e lements in steel promote hardenability. For example , type 4340 steel shown in Fig. 26 has s ignif icant levels o f carbon, manganese , nickel, copper, and molybdenum to promote hardenability. More details about hardenabil i ty can be found in Ref 2.

The mar tens i te start temperature (Ms) for type 4340 is 300 °C (570 °F). Carbon lowers the M s temperature, as shown in Fig. 27, and alloying e lements such as carbon, manganese , chromium, nickel, and mo lybdenum also lower M s temperature. Many empirical equations have been developed over the past 50 years relating M s tempera-

2O0 (271)

Notched impact tests /

"~ (217) ~ e ,~ If energy >:, 120 j ~ ," Transition temperature

(163)

= 80 (106)

~ 40 0 (54) ~ ~ '------

0

~ 100 .cg ~ 80 ~ ~ ' ~ ~'~UIt imate: strength I t ~

= Yield strength 0=

.¢ 60 ~

=m Reduction in area 20 - - Smooth tensile testa I

03

0 0 0.1 0.2 0:3 0.4 0.5 0.6 0.7 0.8 0.9

Carbon, wt%

Fig. 19 Mechanical properties of ferrite-pearlite steels as a function of carbon content. Source: Ref 2

160

120

80

40

-4O

oo

E e

8 ,==

CO ==

o


Temperature, °F

-1 O0 0 100 200 300 400 250 I I I F I

200

¢= • 150 t~ ¢x _E

175

Fig. 20

0.11%C 150

125 a= >~

g 100 f

100 / 0.20% C .." ....................... 0.31 Yo C - 75

_ _ 0 . 4 1 % C 0 6 0 % C -

50 ~~.. . . . . . .~ .""" .- ~ , • ~ " . . . . . . . . ~0.49%C 2 50 ."" • s

, ' ° 1 , , . -" " ' . t . . - " - - . - ' . ' . . . . . . . . . . . oO." •

,.: ... s . o . ° "~ . . " .. " ~ '~ 0.80%C 25

0 • 0

-100 -50 0 50 100 150 200 250

Temperature, °C

Effect Of carbon content in ferrite-peadite steels on Charpy V-notch transition temperature and shelf energy. Sou rce: Ref 17

ture to composit ion. One recent equation by An- drews (Ref 24) is:

M s (°C) = 539 - 423(C) - 30.4(Mn) - 12.1(Cr) - 17.7(Ni) - 7_5(Mo) (Eq 11)

With sufficient alloy content, the M s temperature can be below room temperature, which means that the t ransformat ion is incomplete and retained austenite can be present in the steel.

The microstructure of martensi t ic steels can be generally classed as either lath martensite, plate martensi te , or mixed lath and plate martensite. In plain carbon steels, this classif icat ion is related

to carbon content, as shown in Fig. 27. Lath mar tens i te forms at carbon contents up to about 0.6%, plate martensi te is found at carbon contents greater than 1.0%, and a mixed mar tens i te micros t ructure forms for carbon contents between 0.6 and 1.0%. An example of lath mar ten- site is shown in Fig. 28 and plate martensi te in Fig. 29. Generally, plate martensi te can be distin- guished from lath martensi te by its plate morphology with a central mid-fib. Also, plate mar tens i te may contain numerous microcracks , as shown in Fig. 30. These form during transfor- mat ion when a growing plate impinges on an ex- is t ing plate. Because of these microcracks, plate mar tens i te is generally avoided in most applica-

tions. The important micros t ructura l units measured in lath mar tens i te are lath width and packet size. A packet is a grouping o f laths having a common orientation.

Pla in-carbon and low-al loy martensi t ic steels are rarely used in the as -quenched state because of poor ductility. To increase ductility, these martensi t ic steels are tempered (reheated) to a temperature below 650 °C (1200 °F). During tempering, the carbon that is in supersaturated solid solution precipitates on preferred crystallographic planes (usually the octahedral {111} planes) of the martensi t ic lattice. Because of the preferred orientation, the carbides in a tempered mar tens i te have a character is t ic a r rangement as seen in Fig. 31.

Tempered mar tens i te has s imilar morphological features to type B) (lower) bainite. However, a dist inct ion can be made in terms of the orientation differences of the carbide precipitates. This can be seen by compar ing type B l bainite in Fig. 22 with tempered mar tens i te in Fig. 31. However, unless the carbide morpho logy is observed it is very diff icult to d is t inguish between B] bainite and tempered martensi te .

The hardness of mar tens i te is determined by its carbon content, as shown in Fig. 32. Martensi te attains a m a x i m u m hardness o f 66 HRC at carbon contents of 0.8 to 1.0%. The reason that the hardness does not monotonica l ly increase with carbon is that retained austeni te is found when the carbon content is above about 0.4% (austeni te is much softer than martensi te) . Figure 33 shows the increase in volume percent retained austeni te with increas ing carbon content. Yield strength also increases with increas ing carbon content as seen in Fig. 34. This empir ical relat ionship between the yield s trength and carbon content for un tempered low-carbon mar tens i te is (Ref 25):

YS (MPa) = 413 + 17.2 x 10P(C 1/2) (Eq 12)

Lath mar tens i te packet size also has an inf luence on the yield strength, as shown in Fig. 35. The

(a) (b)

Fig. 21 Microstructure of (a) upper bainite and (b) lower bainite in a Cr-Mo-V rotor steel. 2% nital + 4% picral etch. 500x


(a) (b)

Fig, 22 TEM micmgraphs of (a) upper bainite and (b) lower bainite in a Cr-Mo-V rotor steel

l inear behavior fol lows a Hal l-Petch type rela- t ionship of (d-l/2).

Mos t martcnsi t ic steels are used in the tempered condit ion where the steel is reheated after quenching to a temperature less than the lower critical temperature (Act). Figure 36 shows the decrease in hardness with temper ing temperature for a number o f carbon levels. Plain-carbon or low-al loy martensi t ic steels can be tempered in lower or h igher temperature ranges, depending on the balance of propert ies required. Tempering between 150 and 200 °C (300 and 390 °F) will

mainta in much of the hardness and strength of the quenched martensi te and provide a small im- provement in ductil i ty and toughness (Ref 26). This t reatment can be used for bearings and gears that are subjected to compress ion loading. Tem- pering above 425 °C (796 °F) s ignif icant ly im- proves ductil i ty and toughness but at the expense of hardness and strength. The effect of tempering temperature on the tensile propert ies of a typical o i l -quenched low-alloy steel ( type 4340) is shown in Fig. 37. These data are for a 13.5 mm (0.53 in.) diam rod quenched in oil. The as-

1 0 0 0 ~ 1800 Ac 3 = 930 °C

900 Fs 1 ! 0 1600

.0o ,oo ? - -

o ° 6oo "

5 0 0 - . . . . - - +

~'E ~ l I I " l \ ~ \1 I I \ ~l~""r I I I ~ ~ I I I X I I I II -1800

200 ~ ~ 400

100 200

0 32 10 102 103 104 105

Seconds I ' ' ' ' I ' ' ' ' I ~ ' ' ~ I

1 10 102 103

Minutes I ' I ' ' I ' I 1 4 10 30

Time Hours

Fig. 23 A CCT diagram of a I/2Mo-B steel. Composition: 0.093% C, 0.70% Mn, 0.36% Si, 0.51% Mo, 0.0054% B. Austenitized at Ac 3 + 30 °C for 12 rain. Bs, bainite start; B o bainite finish; Fs, ferrite start; F o ferrite finish. Num-

bers in circles indicate hardness (HV) after cooling to room temperature. Source: Ref 20

quenched rod has a hardness of 601 HB. Note that by temper ing at 650 °C (1200 °F), the hardness (see x-axis) decreased to 293 HB; or to less than hal f the as -quenched hardness . The tensile s t rength has decreased from 1960 MPa (285 ksi) at a 200 °C (400 °F) temper ing temperature to 965 MPa (141 ksi) at a 650 °C (1200 °F) tempering temperature. However, the ductility, represented by total e longat ion and reduction in area, increases dramatically. The tempering process can be retarded by the addition of certain alloying e lements such as vanadium, molybdenum, manganese , chromium, and silicon. Also, for tempering, temperature is mu ch more important than t ime at temperature.

Temper embri t t lement is possible during the temper ing of alloy and low-alloy steels. This em- br i t t lement occurs when quenched-and- tempered steels are heated in, or slow cooled through the 340 to 565 °C (650 to 1050 °F) temperature range. Embri t t lement occurs when the embrit- t l ing elements , ant imony, tin, and phosphorus, concentrate at the austeni te grain boundaries and create intergranular segregat ion that leads to intergranular fracture. The element molybdenum

1200 " ' "q 'e

~-. 1050 ~ ' ~

" 900

750 .== ¢/)

i.~ 600 ~ 1 + Ferrite eel==, peadite

~'~ "/45 450 Martensites Bainites ', ~*

I I I I ~-'- i i 1- 3OO 4OO 500 600 7OO 8O0

Transformation temperature, °C

Fig. 2 4 Relationship between transformation temperature and tensile strength of ferrite-pearlite, bain-

itic, and martensitic steels. Source: Ref 5


900 / t~ IL

to 750

~ QII) •

" 600 • °

450 15 20 25

Grain size (d-1/2), mm -1/2

Fig. 25 Relationship between bainite lath width (grain size) and yield strength. Source: Ref 5

has been shown to be beneficial in prevent ing temper embri t t lement .

The large variation in mechanical propert ies o f quenched-and- tempered martensi t ic steels pro- vides the structural designer with a large number of property combinat ions• Data, like that shown in Fig. 37, are avai lable in the Section "Carbon and Alloy Steels" in this Handbook as well as Volume 1 of the ASM Handbook and the ASM Specialty Handbook: Carbon and Alloy Steels. Hardnesses of quenched-and- tempered steels can be est imated by a method es tabl ished by Grange et al. (Ref 27). The general equation for hardness is:

HV = HV C + AHVMn + AHVp + AHVsi + AHVNi + AHVcr + AHVMo + AHV v (Eq 13)

where HV is the es t imated hardness value (Vick- e r s ) .

In order to use this relat ionship, one m u s t determine the hardness value of carbon (HVc) from Fig. 38. For example , if one assumes that a tempering temperature of 540 °C (1000 °F) is used and the carbon content of the steel is 0.2% C, the HV c value after temper ing will be 180 HV. Sec- ond, the effect o f each al loying e lement mus t be determined from a figure such as Fig. 39. This graph represents a tempering temperature of 540 °C (1000 °F). Graphs represent ing other tempering temperatures can be found in Ref 27.

To illustrate the use of the Grange et al. method, the same type 4340 steel shown in Fig. 37 is used. The composi t ion of the steel is 0.41% C, 0.67% Mn, 0.023% P, 0.018% S, 0.26% Si, 1.77% Ni, 0.78% Cr, and 0.26% Mo. Assuming a 540 °C (1000 °F) tempering temperature, the es- t imated h a r d n e s s value for carbon is 210 HV. From Fig. 38, the hardness values for each of the other alloying e lements are:

Element Ccetent, % Hardaess, HV

Carbon 0.41 210 Manganese 0.67 38 Phosphorus 0.023 7 Silicon 0.26 15 Nickel 1.77 12 Chromium 0.78 43 Molybdenum 0.26 55 Total hardness 380

According to Fig. 37, the hardness value after tempering at 540 °C (1000 °F) was 363 HB (see Brinell hardness values along x-axis) . From the ASTM E 140 convers ion table ( included in the

1600

1400

1200

1000 g

E 800

600

400

200 1 2 5 10 20 50 100 200

Cooling timel s

Fig, 26 The CCT diagram for type 4340 steel austenitized at 845 °C (I 550 °F). Source: Ref 23

870

760

650

540 o ~

425 E

315

205

95 500 1000

examples of hardness conversion tables for steels, which can be found in the Section "Glos- sary of Terms and Engineer ing Data" in this Handbook) , a Brinell hardness of 363 HB equates to a Vickers hardness of 383 HV. The calculated value of 380 HV (in the table above) is very close to the actual measured value of 383 HV. Thus , this method can be used to es t imate a spe- cific hardness value after a quenching-and- tem- pering heat t reatment for a low-al loy steel. Also , as a rough approximation, the derived Brinel l hardness value can be used to es t imate tensi le s t rength by the fol lowing equat ion (calculated f rom ASTM E 140 convers ion table):

TS (MPa) = - 42.3 +3.6 HB (Eq 14)

For the above example , a type 4340 quenched- and- tempered (540 °C, or 1000 °F) steel with a calculated hardness of 363 HB would have an es t imated tensile s t rength f rom Eq 14 of 1265 MPa (183 ksi). From Table 1, this measured tensile s t rength of a type 4340 quenched-and- tem- pered (540 °C, or 1000 °F) steel is 1255 MPa (182 ksi).

It is seen that quenched-and- tempered martensitic steels provide a wide range of properties. The design engineer can choose from a large number of pla in-carbon and low-al loy steels. In

870 1600 .8o 650

o ~ cff 540

425 ~""~. , ,~

315

205

95 l% ~: ~ . . . . . . . . .

Lath .... ?~ Mixea / , ~ ~:,: D20

0 0.2 0.4 0.6 0.6

1400

Carbon, wt%

1200

i i o lOOO =d

800 Q_

E

~o ff

Fig. 27 Effect of carbon content on M s temperature in steels. Source: Ref 6

400

Plate 0

1.0 1.2 1.4 1.6


Fig. 28 Microstructure of a typical lath martensite. 4% picral + HCI. 200x Fig. 29 Microstructure of a typical plate martensite. 4% picral + HCI. 1000x

addit ion to this large list of steels, there are two other commercia l ly important categories o f fully mar tensi t ic steels, namely, martensi t ic s ta inless steels and maraging steels.

Like the ferritic s ta inless steels, martensi t ic s ta inless steels (e.g., type 403, 410, 414, 416, 420, 422, 431, and 440) are h igh -ch romium iron al loys (12 to 18% Cr), but with deliberate addi- t ions of carbon (0.12 to 1.2% C). These steels use carbon in order to stabilize austeni te in iron- ch romium alloys (Fig. 12). The expanded region of austeni te is called the y-loop. In the Fe-Cr phase d iagram (without C), the y-loop extends to about 12% Cr (see Fig. 12). With carbon addi- t ions, austeni te can exist up to 25% Cr. T h e s e steels can be heat treated much like those of the low-al loy steels. However, martensi t ic s ta inless steels , with such h igh ch romium contents , can form martensi te on air cooling, even in thick sections. Martensi t ic s ta inless steels are considered

h igh-s t rength stainless steels because they can be treated to achieve a yield s t rength between 550 MPa (80 ksi) and 1725 MPa (250 ksi), as seen in Table 1. On the other hand, ferritic s tainless steels, which do not contain carbon, are not considered h igh-s t rength steels because their yield s trength range is only 170 to 450 MPa (25 to 64 ksi). Because of their high s t rength and hardness , coupled with corrosion resis tance, martensi t ic s ta inless steels are used for knives and other appl icat ions requiring a cut t ing edge as well as some tool steel applications (for example, molds for producing plastic parts).

Maraging steels are a separate class of martensitic steels and are considered ul t rahigh-s t rength steels with yield strength levels as high as 2500 MPa (360 ksi), as seen in Table 1. In addition to extremely h igh strength, the maraging steels have excel lent ductility and toughness. These very- low carbon steels contain 17.5 to 18% Ni,

8.5 to 12.5% Co, 4 to 5% Mo, 0.20 to 1.8% Ti, and 0,10 to 0.15% A1. Because of the high alloy content, especial ly the cobalt addition, they are very expensive. Their high s trength is developed by austeni t iz ing at 850 °C (1560 °F), fol lowed by air cool ing to room temperature to form lath martensi te . However, the martensi t ic consti tuent in maraging steels is relatively so f t - -28 to 35 H R C - - w h i c h is an advantage because the component can be machined to final form directly upon cooling. The final s tage o f s t rengthening is through an aging process, carried out at 480 °C (900 °F) for 3 h. During aging, the hardness increases to about 51 to 58 HRC depending on the grade o f maraging steel. The aging t reatment pro- motes the precipitat ion of a rodlike intermetall ic compound Ni3Mo. These precipitates can only be observed at high magnif icat ion (e.g., by TEM). The precipitates s t rengthen the surrounding matrix as they form during aging. Full hardening

Fig. 30 Microcracks formed in plate martensite. 4% picral + HCl/sodium metabisulfite Fig. 31 Ttransmission electron micrograph showing carbide morphology in tempered etch. 1000x martensite


> "I-

¢=

t~ "1-

900

800

700

600

500

400

300

200

100 0

65

6o

o t~ -.r of

so ~

40

30

20 1o 0

J I I I I I 0.2 0.4 0.6 0.8 1.0 1.2

Carbon, wt%

Fig. 32 Effect of carbon content on the hardness of martensite. Source: Ref 4

can be developed, even in very thick sections. Maraging steels are used for die-casting molds and aluminum hot-forging dies as well as numerous aircraft and missile components.

Austenite

Austenite does not exist at room temperature in plain-carbon and low-alloy steels, other than as small amounts of retained austenite that did not transform during rapid cooling. However, in certain high-alloy steels, such as the austenitic stainless steels and Hadfield austenitic manganese steel, austenite is the microstructure. In these steels, sufficient quantifies of alloying elements that stabilize austenite at room temperature are present (e.g., manganese and nickel). The crystal structure of austenite is face-centered cubic (fee) as compared to ferrite, which has a (bcc) lattice. A fcc alloy has certain desirable characteristics; for example, it has low-temperature toughness, excellent weldability, and is non- magnetic. Because of their high alloy content, austenitic steels are usually corrosion resistant. Disadvantages are their expense (because of the

260

- 220 J~

180 "o

"~ 140

~. 100 o

6 0

o o -~,~1 1700

t300 -lltOO

o,° "o

so "~

Y t oo I I I I

0 0.10 0.20 0.30 0.40

Carbon content, wt%

Fig. 34 Relationship between carbon content and the yield strength of martensite. Source: Ref 4

>o

• ->¢ 100 =o

"~ 75 8 t~ 03 E

25

L ] M s temperature

Lath marten ' l i t e , ' ~ ~ ~ "~1 . . . . . . . . . . ~ - - relative volh°/° - - -~ ~ ~ /

Retained 7, vol% . t ~ I I I 0.4 0.8 1.2 1.6

Carbon, wt%

700

500

3OO

100

<

40

20

o

E

o p -

Fig. 33 Effect of carbon content on the volume percent of retained austenite (7) in as-quenched martensite. Source: Ref 4

alloying elements), their susceptibility to stress- corrosion cracking (certain austenitic steels), their relatively low yield strength, and the fact that they cannot be strengthened other than by cold working, interstitial solid-solution strengthening, or precipitation hardening.

The austenitic stainless steels (e.g., type 301, 302, 303, 304, 305,308, 309, 310, 314, 316, 317, 321, 330, 347, 348, and 384) generally contain from 6 to 22% Ni to stabilize the austenite at room temperature. They also contain other alloying elements, such as chromium (16 to 26%) for corrosion resistance, and smaller amounts of manganese and molybdenum. The widely used type 304 stainless steel contains 18 to 20% Cr and 8 to 10.5% Ni and is also called 18-8 stainless steel. From Table 1, the yield strength of annealed type 304 stainless steel is 290 MPa (40 ksi), with a tensile strength of about 580 MPa (84 ksi). However, both yield and tensile strength can be substantially increased by cold working as shown in Fig. 40 (see Table 1). However, the increase in strength is offset by a substantial decrease in ductility, for example, from about 55%

elongation in the annealed condition to about 25% elongation after cold working.

Some austenitic stainless steels (type 200, 201, 202, and 205) employ interstitial solid-solution strengthening with nitrogen addition. Austenite, like ferrite, can be strengthened by interstitial elements such as carbon and nitrogen. However, carbon is usually excluded because of the deleterious effect associated with precipitation of chromium carbides on austenite grain boundaries (a process called sensitization). These chromium carbides deplete the grain-boundary regions of chromium, and the denuded boundaries are extremely susceptible to corrosion. Such steels can be desensitized by heating to high temperature to dissolve the carbides and place the chromium back into solution in the austenite. Nitrogen, on the other hand, is soluble in austenite and is added for strengthening. To prevent nitrogen from forming deleterious nitrides, manganese is added to lower the activity of nitrogen in the austenite, as well as to stabilize the austenite. For example, type 201 stainless steel has composition ranges of 5.5 to 7.5% Mn, 16 to 18% Cr,

ASTM grain size 2 4 6 8 10 12

1200 i i = i i = 120

a. E ,50

"~ 800 o-o~ 80

600 ~ ~ -r.,

400 ~ . ~ 40 o

0 2 4 6 8 10 12 14 16

Lath martensite packet size (d-1/2), mm -1/2

Fig. 35 Relationship between lath martensite packet size (dl and yield strength of Fe-0.2%C (upper line) and Fe-Mn (lower line) martensites. Source: Ref 2


70

Tempering temperature,°F

200 400 600 800 1000 1200 1400

6 0

0 ,0,,o ~ - o . ~ c %

n-° 50 ~ ~ ~ ~ o ~

= 0=00=0; "l- ~ 40 ~ m = ~ ~ 0.10"0.20 ~/o C ~ ~

30

20

\ 1 0 I I I I I I

As- 100 200 300 400 500 600 700

quenched Tempering temperature, °C

Fig. 36 Decrease in the hardness of martensite with tempering temperature for various carbon contents. Source: Ref 2

3.5 to 5.5% Ni, and 0.25% N. The other type 2xx series of steels contain from 0.25 to 0.40% N.

Another important austenitic steel is austenitic manganese steel. Developed by Sir Robert Had- field in the late 1890s, these steels remain austenitic after water quenching and have consid- erable strength and toughness. A typical Hadfield manganese steel will contain 10 to 14% Mn, 0.95 to 1.4% C, and 0.3 to 1% Si. Solution annealing is necessary to suppress the formation of iron carbides. The carbon must be in solid solution to stabilize the austenite. When completely austeni-

900 As-quenched

800 65 As-quenched / ~0~°F I

hardiness " ~ - / / I 400°F '= / ° °

0= ], / . . , ° " / b ' / ' ' = "f" 600 °F..,@ n- ,6 z~" l " - ~.c. 50 ="

"O i l 45

400 800 °F"°" 40 ~ • -. ~,,,,, ~o- I

300 ; ; ~ o ~ ~. , , ,~ ~ , - , 1000 OF .o" 30

200 ~....,~ 1200 °F=o= 1 3 0 0 °F

100 I 0.2 0.4 0.6 0.8 1.0

Carbon, %

Fig. 38 Relationship between hardness of tempered martensite with carbon content at various tem-

pering temperatures. Source: Ref 2

I1.

J::

o}

Hardness, HB 555 477 415 363

300 (2070)

250 (172o)

200 (13oo)

150 (1030)

100 (690)

293

\ \

N Tensile strength

' ' ' \ N - - Yield point - N N

N i X

- - Reduction in area --~

Elongation ' i ~ '

400 600 800 1000 1200 (200) (320) (430) (540) (650)

Tempering temperature, °F(°C)

¢5 t~

70 6O

4O

!i 5 ILl

Fig. 37 Effect of tempering temperature on the mechanical properties of type 4340 steel. Source:

Ref 2

tic, these steels can be work hardened to provide higher hardness and wear resistance. A work- hardened Hadfield manganese steel has excellent resistance to abrasive wear under heavy loading. Because of this characteristic, these steels are ideal for jaw crushers and other crushing and grinding components in the mining industry. Also, Hadfield manganese steels have long been used for railway frogs (components used at the junction point of two railroad lines).

Ferrite-Cementite

When plain-carbon steels are heated to temperatures just below the lower critical tempera-

ture (Ac]), the process of spheroidization takes place. Figure 41 shows a fully spheroidized steel microstructure. The microstructure before spheroidization is pearlite. During spheroidization, the cementite lamellae of the pearlite must change morphology to form spheroids. The process is controlled by the diffusion rate of carbon and portions of the lamellae must "pinch-off" (dissolve) and that dissolved carbon must diffuse to form a spheroid from the remaining portions of lamellae. This process takes several hours. Spheroidization takes place in less time when the starting microstructure is martensite or tempered martensite. In this process, the spheroidized carbides are formed by growth of carbides formed during tempering.

A fully spheroidized structure leads to improved machinability. A steel in its fully sphe-

8O

70

> 60 "1-

~ 50

~ 4o .E

~ 3o

o 20

lO

• Mo

/ e •

• tw p

j .S/ !° Cr

o Si

f j

0 0.02 0.04 o.o6 0.1 0.2 0.4 0.6 1 2

Element content, %

Fig. 39 E ~ of alloying elements on the retardation or softening during tempering at 540 °C (1 000 °F) relative to iron- carbon alloys. Source: Ref 2

Strudure/Property Relationships in Irons and Steels / 171

1200 /

100o / f -- Tensile/~stStrength - - /

~ mo I rength

= / 600 / ' (0.2*/. offset)

~- 40 =m o 200 ~- []

~ 20

0 ~ 0 10 20 30 40 50 60

Cold work. %

Fig. 40 Influence of cold work on mechanical properties of type 304 stainless steel. Source: Ref 4

roidized state is in its softest possible condition. Some steels, such as type 1020, are spheroidized before cold forming into tubing because spheroidized steels have excellent formability.

Ordinary low-carbon, cold-rolled, and annealed sheet steels have ferritic microstructures with a small amount of grain-boundary cementite, as shown in Fig. 8. These carbides nucleate and grow on the ferrite grain boundaries during the annealing process, which takes place in the lower portion of the intercritieal temperature region (i.e,, the region between the A 3 and A 1 temperatures shown in the iron-carbon diagram, Fig. 6). Many modern-day automotive sheet steels are produced with very low carbon levels to avoid these grain-boundary carbides because they de- grade formability.

Ferrite-Martensite

A relatively new family of steels called dual- phase steels consists of a microstracture of about

Fig. 42 Microstructure of a typical dual-phase steel. 2% nital etch. 250x

0 ~+, ,+ , . . + .,yff/S+~ ++ z9 O,o,:" Q ~ - * o "

s ~ t " ~ u t z , o v # = = o © " x ~, o+ DO . ",-,~', . o o ~ o • ~ , .o0. ° % 0 : ' : , ~ , - , , ~ . o- ~

• . " , " ~ . t ) •

OC~ 06 o +' )1 * . ,~ ~ ====, ¢'~'~ o " ( / (7 ,,,,. I=, .',,a 0 o o o oo v . ,0 I= ~F' = ,~,=p " %

O . . + + o , , o 0+,o+ . ~ l . . ~ , , o . : .o ~ , . = . . . . ~ + • ~ o +° "o. o ~>/2 + .~.,o,'-." • , . "1 " . ~ == ~+)1 ^+L

~:~ , e - ~ a ~, ~" 0,//~,,.. ,,- .. oOo..~:" 0~"o o:.- "'+,~" +o~"~'~ ~o'+<3 c o 2 + : + " ' . oo ,,? a : - - , £ . 0- v++o

• . ¢ ~ ' , , * " o o ~ '

°e:+ e" ooV+ '~ d'-" o~;~, + " \

Fig, 41 Microstructure of a fully spheroidized steel. 4% picral etch. 1000x

15 to 20% martensite in a matrix of ferrite. The microstructure of a typical dual-phase steel is shown in Fig. 42. In most plain-carbon and low- alloy steels, the presence of martensite in the microstructure is normally avoided because of the deleterious effect that martensite has on ductility and toughness. However, when the martensite is embedded in a matrix of ferrite, it imparts desirable characteristics. One desirable characteristic is that dual-phase steels do not exhibit a yield point. Figure 43 compares the stress-strain behavior of four steels: plain carbon, SAE 950X, and SAE 980X, which exhibit a yield point with the fourth, a dual-phase steel (GM 980X). This means that the cosmetically unappealing Ltiders bands that form during the discontinuous yielding (i.e., yield point) are absent in a dual-phase steel. Also note in Fig. 43 that the dual-phase steel has much more elongation than the SAE

980X of similar tensile strength. These characteristics are especially important in formability.

A unique characteristic of a ferrite-martensite dual-phase steel is its substantial work hardening capacity. This allows the steel to strengthen while being deformed. By proper design of the stamping dies, this behavior can be exploited to produce a high-strength component. Most conventional high-strength steels have limited formability because their high strength is developed prior to the forming process.

FerriteoAustenite

High-alloy steels having approximately equal proportions of fcc austenite and bcc ferrite, with ferrite comprising the matrix, are referred to as duplex stainless steels. The microstructure of a typical duplex stainless steel is shown in Fig. 44. Although the exact amount of each phase is a function of composition and heat treatment, most alloys are designed to contain about equal amounts of each phase in the annealed condition.

100

~" (552)(690)80 ~ / . . . . . . . . ~',1 'IGM 980X'

~v 60 t" ~ J ~ / (414) ~ ' ~

/ SAE 950X 40

~ (276) / Plain c a / ~ n ~

I 038) 0

0 10 20 30 40 Strain in two-inch gage length, %

Fig. 43 Comparison of the stress-strain curves of three discontinuously yielding sheet steels (plain car-

loon, 5AE 950X, and SAE 980X) and a dual-phase steel (GM 980X). in addition to the differences in yielding behavior, note the higher percentage of uniform elongation in the dual-phase steel compared with the conventional SAE 980X of similar tensile strength. Source: Ref 2


The duplex structure resul ts in improved stress- corrosion cracking resis tance, compared with austeni t ic s ta inless steels, and improved toughness and ductility, compared with the ferritic s ta inless steels. Duplex s ta inless steels are capa- ble of tensi le yield s t rengths ranging f rom 400 to 550 MPa (60 to 80 ksi) in the annealed condit ion, which is approximate ly twice the s t rength of ei ther phase alone.

The principal al loying e lements in duplex s ta inless steels are ch romium and nickel, but nitrogen, molybdenum, copper, si l icon, and tung- sten may be added to control s tructural balance and to impar t certain corros ion-res is tance characteristics. Four commercia l groups o f duplex s ta inless steels, listed in order o f increas ing corrosion resis tance, are:

Fe-23Cr-4Ni-0.1N • Fe-22Cr-5 .5Ni-3Mo-0.15N • Fe-25Cr-5Ni-2 .5Mo-0 .17N-Cu • Fe-25Cr-7Ni-3 .5Mo-0 .25N-W-Cu

Because o f their excel lent corrosion resis tance, ferr i te-austenite duplex s ta inless steels have found widespread use in a range o f industr ies , particularly the oil and gas, petrochemical , pulp and paper, and pollut ion control industr ies . They are commonly used in aqueous, chlor ide-containing env i ronments and as rep lacements for austeni t ic s ta inless steels that have suffered s t ress-corrosion cracking or pit t ing dur ing ser- vice.

Graphite When carbon contents o f i ron-carbon alloys

exceed about 2%, there is a tendency for graphite to form (see Fe-C diagram in Fig. 6b). This is especial ly true in gray cast iron in which graphite f lakes are a predominant microstructural fea- ture (Fig. 3). Gray cast iron has been used for centuries because it mel ts at a lower temperature than steel and is easy to cast into various shapes. Also, the graphite flakes impar t good machinability, act ing as chip breakers, and they also provide excel lent damping capacity. Damping capacity is important in mach ines that are subject to vibration. However, gray cast iron is l imited to appl icat ions that do not require toughness or ductility, for example , total e longat ion o f less than 1%. The flake morphology of the graphi te pro- vides for easy crack propagat ion under applied stress.

Gray cast irons usual ly contain 2.5 to 4% C, 1 to 3% Si, and 0.1 to 1.2% Mn. The graphite flakes can be present in five different morpholo-

Type A Type B

./T.¢,i!

, / I/ / ; [_,~ '~"

72,: • / ,'~..,,~, ,, ,,. -- '~,1 '~. ' •

Uniform distribution, Rosette grouping, random orientation random orientation

Fig. 45 Classification of different graphite flake morphology

Fig. 44 Microstructure of a typical mill-annealed duplex stainless steel plate showing elongated austenite islands in the ferrite matrix. Etched in 15 mL HCI in 100 mL ethyl alcohol. 200x

gies as seen in Fig. 45. Type A, because o f its r andom orientation and distribution, is preferred in many applications, for example, cyl inders of internal combust ion engines. The matr ix of a typical gray cast iron is usual ly pearlite. How- ever, ferrite-pearlite or martensi t ic micro- s t ructures can be developed by special heat treat- ments . As a structural material, gray cast i ron is selected for its high compress ive strength, which ranges f rom 572 to 1293 MPa (83 to 188 ksi), a l though tensile s trengths of gray iron range only from 152 to 431 MPa (22 to 63 ksi). Gray cast i rons are used in a wide variety of applicat ions, inc luding automot ive cyl inder blocks, cyl inder heads and brake drums, ingot molds, machine hous ings , pipe, pipe fi t t ings, manifolds , compres- sors, and pumps.

Another form of graphite in cast iron is spheroidal graphite found in ductile cast irons (also called nodular cast irons). The microstructure of a typical ducti le cast iron is shown in Fig. 46. This form of graphite is produced by a process called inoculation, in which a magne- s ium or cerium alloy is thrust into mol ten cast iron immediate ly prior to the cast ing operation. These e lements form intermetall ic compounds that act as a nucleat ing surface for graphite. With a spherical morphology, the graphite no longer renders the cast iron brittle as do graphite f lakes in gray cast iron. Ducti le irons have much higher ducti l i ty and toughness than gray iron and thus expand the use of this type of ferrous alloy. Mos t ducti le iron cast ings are used in the as-cast form. However, heat t reatment can be employed to al-

ter the matr ix microstructure to obtain desired properties. The matrix can be fully ferritic, fully pearli t ic, fully martensi t ic , or fully bainitie, depending on composi t ion and heat treatment. The yield s trength o f typical ductile cast irons ranges f rom 276 to 621 MPa (46 to 76 ksi), and their tensi le s t rengths range from 414 to 827 MPa (60 to 120 ksi). Total e longat ion ranges f rom about 3 to 18%. Heat treated, aus tempered ductile irons have yield s t rengths ranging f rom 505 to 950 MPa (80 to 138 ksi), tensi le s trengths ranging from 860 to 1200 MPa (125 to 174 ksi), and total e longat ions ranging from 1 to 10%. Uses for due- tile iron include gears, crankshaf ts , paper-mill dryer rolls, valve and pump bodies, steering knuckles , rocker arms, and var ious machine components .

Cementite

A major microstructural const i tuent in white cast i ron is cementi te . The microstructure of a typical white cast iron is shown in Fig. 47. The cement i te forms by a eutectic reaction during so- l idification:

Liquid ~-~ Cementite + Austeuite (Eq 15)

The eu tec t i c cons t i tuen t in whi te cast iron is cal led ledeburite and has a two-phase morphology shown as the smal ler part icles in the white matrix in Fig. 48. The eutectic is shown in the Fe-C

Type C

Superimposed flake size, random orientation

Type D

.~>..'

• i:t"

~,~.<.'.,..~ . ~ ~.:

Interdendritic segregation, random orientation

Type E

Interdendritie segregation, preferred orientation


Fig, 46 Microstructure of a typical ductile (nodular) cast iron showing graphite in the form of spheroids.

2% nital etch. 200x. Courtesy of A.O. Benscoter, Lehigh University

binary diagram in Fig. 6(b). The austenite in the eutectic (as well as the austenite in the primary phase) transforms to pearlite, ferrite-pearlite, or martensite, depending on cooling rate and composition. Because of the high percentages of cementite, white cast irons are used in applications requiring excellent wear and abrasion resistance. These irons contain high levels of silicon, chromium, nickel, and molybdenum and are termed alloy cast irons. Such applications include steel mill rolls, grinding mills, and jaw crushers for the mining industry. Hardness is the primary mechanical property of white cast iron and ranges from 321 to 400 HB for pearlitic white iron and 400 to 800 HB for alloy (martensitic) white irons.

REFERENCES

1. P.D. Harvey, Ed., Engineering Properties of Steel, American Society for Metals, 1982

2. G. Krauss, Principles of the Heat Treatment of Steel, American Society for Metals, 1980

3. R.W.K. Honeycombe, Steels--Microstructure and Properties, American Society for Metals, 1982

4. W.C. Leslie, The Physical Metallurgy of Steels, McGraw-Hill, 1981

5. F.B. Picketing, Physical Metallurgy and the Design of Steels, Applied Science, 1978

6. G. Krauss, Microstruetures, Processing, and Properties of Steels, Properties and Selection: Irons, Steels, and High-Performance Alloys, Vol 1, ASM Handbook, 1990, p 126

7. E.C. Bain and H.W. Paxton, Alloying Elements in Steel, 2nd ed., American Society for Metals, 1961, p 62

8. Microalloying 75, Conference Proceedings (Washing- ton, D.C., Oct 1975), Union Carbide Corp., 1977, p 5

9. T.B. Massalski, J.L. Murray, L.H. Bennett, and H. Baker, Ed., Binary Alloy Phase Diagrams, Vol 1, American Society for Metals, 1986, p 822

10. W. Haller, 1L Schweitzer, and L. Weber, Can. Metall. Q., Vol 21 (No. 1), 1982, p 3

11. J.M. Hyzak and I.M. Bernstein, Metall. Trans. A, Vol 7A, 1976, p 1217

Fig, 47 Microstructure of a typical white cast iron. 4% picral etch. 100x. Courtesy of A.O. Benscoter, Lehigh University

Fig. 48 Microstructure of the eutectic constituent ledebutite in a typical white cast iron. 4% picral etch. 500x. Courtesy of A.O. Benscoter, Lehigh University

12. G.E Vander Voort and A. Ro6sz, Metallography, Vo117 (No. 1), 1984, p 1

13. H. Iehinose et al., paper 1.3, Proc. First Int. Heavy Hauls Railway Conf., Association of American Rail- roads, 1978, p 1

14. G.E Vander Voort, Ed.,Atlas of Time-Temperature Dia- grams for Irons and Steels, ASM International, 1991, p 570

15. B.L. Bramfitt, Proc. 32nd Mechanical Working and Steel Processing Conference, Vol 28, ISS-AIME, 1990, p 485

16. F.B. Picketing, Towards Improved Toughness and Duc- tility, Climax Molybdenum Co., 1971, p 9

17. G.J. Roe and B.L. Bramfitt, Notch Toughness of Steels, Properties and Selection: Irons, Steels, and High-Per- formance Alloys, Vol 1, ASM Handbook, ASM Interna- tional, 1990, p 739

18. E.C. Bain, The Sorby Centennial Symposium on the History of Metallurgy, TMS-AIME, 1963, p 121

19. B.L. Bramfitt and J.G. Spoer, MetalL Trans. A, Vol 21A, 1990, p 817

20. G.E Vander Voort, Ed., Atlas of Time-Temperature Dia- grams for Irons and Steels, ASM International, 1991, p 249

21. W. Steven and A.G. Haynes, J. Iron Steel Inst., Vol 183, 1956, p 349

22. R.W.K. Honeycombe and F.B. Pickering, Metall. Trans. A, Vol 3A, 1972, p 1099

23. G.E Vander Voort, Ed., Atlas of~me-Temperature Dia- grams for Irons and Steels, ASM Intgrantional, 1991, p 544

24. K.W. Andrews, J. Iron Steel Inst., Vol 203, 1965, p 271 25. G.R. Speich and H. Warlimont, J. Iron Steel Inst., Vol

206, 1968, p 385 26. G. Kranss, Z Iron Steel Inst. Jpn., Int., Vol 35 (No. 4),

1995, p 349 27. R.A. Grange, C.R. Hrihal, and C.F. Porter, Metall.

Trans. A, Vol 8A, 1977, p 1775

Structure/Property Relationships in Irons and Steels

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